Alloy flake for rare earth magnet, production method thereof, alloy powder for rare earth sintered magnet, rare earth sintered magnet, alloy powder for bonded magnet and bonded magnet

ABSTRACT

Disclosed is a rare earth magnet in the R-T-B (rare earth element-transition metal-boron) system that is made from an improved composition and properties of main phase alloy in the R-T-B system containing Pr and a boundary alloy. Disclosed also is a manufacturing method of the rare earth magnet alloy flake by a strip casting method with improved rotating rollers such that the alloy flake has a specified fine surface roughness and has a small and regulated amount of fine R-rich phase regions. Consequently, the alloy flake for the rare earth magnet does not containing α-Fe and has a homogeneous morphology so that the rare earth magnet formed by sintering or bonding the alloy flakes exhibits excellent magnetic properties.

CROSS-REFERENCE TO RELATED APPLICATION

This application claims the benefit pursuant to 35 U.S.C. §119(e)(1) ofU.S. Provisional Applications, No. 60/343,187 filed on Dec. 31, 2001,U.S. Provisional Application No. 60/343,192 filed on Dec. 31, 2001, U.S.Provisional application No. 60/410,802 filed on Sep. 16, 2002, and aU.S. Provisional Application No. 60/430,649 filed on Dec. 4, 2002.

TECHNICAL FIELD

The present invention relates to a main phase alloy containing Pr and aboundary phase alloy for producing a rare earth magnet, to a method forproducing the alloy, to a mixed powder for a rare earth sintered magnet,for a rare earth magnet; and to a rare earth magnet. The presentinvention also relates to rare earth magnet alloy flake, formed of anR-T-B alloy (R represents at least one rare earth element including Y; Trepresents transition metals including Fe as an essential element; and Brepresents boron); a method for producing the flake; to a rare earthsintered magnet alloy powder, a rare earth sintered magnet, a bondedmagnet alloy powder and a bonded magnet, and more particularly, relatesto a rare-earth-containing alloy flake produced through the stripcasting method and to a method for producing the alloy flake.

BACKGROUND ART

In recent years, production of Nd—Fe—B alloys serving as magnet alloyshas sharply increased by virtue of excellent characteristics of thealloys, and these alloys are employed in HDs (hard disks), MRI (magneticresonance imaging), a variety of motors, etc. Typically, Nd is partiallysubstituted by another rare earth element such as Pr or Dy and Fe ispartially substituted by another transition metal element such as Co orNi. Such substituted alloys as well as Nd—Fe—B alloys are generallyreferred to as R-T-B alloys. Herein, R represents at least one rareearth element including Y, and T represents transition metals includingFe as an essential element. Fe may be partially substituted by Co or Ni.Other elements such as Cu, Al, Ti, V, Cr, Mn, Nb, Ta, Mo, W, Ca, Sn, Zr,and Hf may be added, to the R-T-B alloys, singly or in combination oftwo or more species. B represents boron and B may be partiallysubstituted by C or N.

An R-T-B alloy contains, as the main phase, a ferromagnetic phase formedof R₂T₁₄B crystals, which contribute to magnetization, and a nonmagneticR-rich phase having a low melting point and containing a rare earthelement(s) at high concentration. Since the R-T-B alloy is an activemetallic material, the alloy is generally melted and mold-cast in vacuumor under inert gas. In a typical method of producing a sintered magnet,an ingot of the alloy is pulverized to powder having a particle size ofabout 3 μm (as measured by means of FSSS (Fisher Sub-Sieve Sizer)); thepowder is subjected to press-forming in a magnetic field; the resultantcompact is sintered in a sintering furnace at a temperature as high asapproximately 1,000 to 1,100° C.; and in accordance with needs, thesintered product is heated, mechanically processed, and plated forcorrosion prevention.

The R-rich phase plays the following important roles in sintered magnetsformed of the R-T-B alloy.

-   (1) Since the R-rich phase has a low melting point, the phase    liquefies during sintering, thereby contributing to achievement of    high density of the resultant magnet, leading to improved    magnetization.-   (2) The R-rich phase functions to smoothen grain boundaries, thereby    reducing the number of nucleation sites of reversed magnetic    domains, thereby enhancing the coercivity.-   (3) The R-rich phase magnetically insulates the main phase, thereby    enhancing the coercivity.

The distribution of the R-rich phase in a magnet, the final product,depends greatly on the microstructure of the as-cast alloy ingot.Specifically, when the alloy ingot is produced through mold casting, aslow cooling rate of the cast ingot often results in formation of largecrystal grains in the R₂T₁₄B phase, and R-rich phase forms largeaggregates which are locally present in the ingot. Thus, the particlesize of the pulverized alloy ingot becomes considerably smaller than thegrain size of crystals present in the R₂T₁₄B phase or than the size ofdispersed R-rich phase. Therefore, particles formed only of the mainphase (R₂T₁₄B phase) containing no R-rich phase and particles formedonly of the R-rich phase are produced, whereby homogeneously mixing themain phase and R-rich phase becomes difficult.

Another problem involved in mold casting is that γ-Fe tends to be formedas primary crystals, due to the slow cooling rate. At approximately 910°C. or lower, γ-Fe transforms into α-Fe, which deteriorates pulverizationefficiency during production of sintered magnets. If α-Fe remains evenafter sintering, magnetic characteristics of the sintered product aredeteriorated. Thus, the ingot obtained through mold casting must besubjected to homogenization treatment at high temperature for a longperiod of time in order to remove α-Fe.

In order to solve the above problems; i.e., segregation of R-rich phaseand precipitation of α-Fe, the strip casting method (abbreviated as SCmethod) and the centrifugal casting method (abbreviated as CC method),which ensure a cooling rate during casting of R-T-B alloy that is fasterthan that attainable by mold casting, are proposed and employed inactual production steps.

In the SC method, molten alloy is poured onto a rotatable copper rollerfor casting, the inside of which is cooled by water, and is formed intoa strip having a thickness of about 0.1 to about 1 mm. During casting,the molten alloy is solidified through rapid cooling, to thereby preventprecipitation of α-Fe which is formed during mold casting and yield analloy having a microcrystalline structure in which R-rich phase isminutely dispersed. Since the R-rich phase is minutely dispersed in thealloy ingot produced through the SC method, dispersion of R-rich phasein the product obtained by pulverizing and sintering the alloy alsobecomes satisfactory, to thereby successfully produce magnets ofimproved magnetic characteristics (Japanese Patent Application Laid-Open(kokai) Nos. 5-222488 and 5-295490).

Meanwhile, the CC method includes feeding a molten metal into theinterior of a cylindrical mold which is rotating, to therebysimultaneously deposit and solidify the molten metal. Thus, the methodattains an intermediate solidification rate between that attainable bymold casting and that attainable by the SC method (U.S. Pat. No.2,817,624). This particular solidification condition is confirmed to beeffective for producing a boundary phase alloy for employment in thetwo-alloy blending method (U.S. Pat. No. 3,234,741).

As compared with the mold casting method, the SC method and CC methodattain high uniformity in microstructure. The uniformity inmicrostructure can be evaluated in terms of crystal grain size anddispersion state of R-rich phase as well as presence of precipitatedα-Fe. When an alloy is cast through the mold casting method, a portionof the resultant alloy ingot which has remained in the vicinity of themold and has been rapidly cooled exhibits a microstructure formed ofminute equiaxed crystal grains called chill crystals and contains acomparatively finely dispersed R-rich phase. However, in the centerportion of the alloy ingot where solidification is finally complete,crystal grains have a large grain size and R-rich phase forms aggregatesin some regions, because of a considerably slow solidification rate inthe center portion.

In the alloy flakes produced through the SC method, chill crystals maybe formed on the side which has been in contact with a rotating rollerfor casting (hereinafter referred to as a mold side). However, anappropriately minute and uniform microstructure can be generallyobtained through rapid-cooling solidification. In addition, sinceformation of α-Fe is suppressed, uniformity in R-rich phase contained ina sintered magnet, a final product, is enhanced, thereby preventingimpairment of α-Fe in terms of pulverizability and magneticcharacteristics.

When a molten alloy is cast through the CC method, the molten alloy isgradually deposited and the thus-solidified thin layers are stacked.Therefore, the cast product can possess a microstructure which is almostuniform from the mold side to the free surface side, despite its largethickness, except that chill crystals are formed in a mold side portion.However, since a conventional CC method (e.g., a method disclosed inU.S. Pat. No. 2,817,624) employs a comparatively high rate of feedingmolten alloy, the substantial solidification rate becomes slower thanthat employed in the SC method. Thus, the conventional CC method attainsan effect for preventing precipitation of α-Fe to a degree, which liesbetween that attainable by mold casting and that attainable by the SCmethod.

In recent years, Nd in R-T-B alloys for producing rare earth magnets hasoften been partially substituted by Pr. This is because partialsubstitution of Nd by Pr causes only a small variation incharacteristics; Pr is less expensive than Nd; and production cost canbe reduced. In the case of an R₂Fe₁₄B compound, saturation magnetizationat room temperature of the compound (R═Nd) is known to be approximately4% higher than that of the compound (R═Pr), but anisotropic magneticfield of the compound (R═Pr) is known to be approximately 5% higher thanthat of the compound (R═Nd). Regardless of whether R is Nd or Pr, phaseconditions around R₂Fe₁₄B compound are substantially the same. Thus,even when Nd of R₂Fe₁₄B is partially substituted by Pr, phaseconstitutions remain substantially unchanged, and magnetism is notdeteriorated by such a subtle change in microstructure.

The present invention is constituted by four aspects and respectiveaspects have following problems to be solved.

The problems to be solved by the first aspect of the present inventionis described below.

From the viewpoint of cost reduction and effective utilization ofresources, substitution in terms of R in R-T-B alloys for producing rareearth magnets, i.e., partial substitution of Nd by Pr, has been widelyemployed. However, the Pr content of R can be elevated up to about 10%by mass, because Pr is chemically active as compared with Nd. Such highchemical activity causes problematic oxidation during production ofmagnets or in the produced magnets.

As compared with the single-alloy method, the two-alloy blending method,which is widely employed for producing high-performance magnets, imposesmore severe limitation on the amount of Pr added. The two-alloy blendingmethod employs two types of raw material alloys; i.e., a main phasealloy, which predominantly provides R₂Fe₁₄B phase (main phase) and has acomposition similar to that of R₂Fe₁₄B, and a boundary phase alloy,which predominantly provides R-rich phase (grain boundary phase) and hasa TRE (Total Rare Earth content) greater than that of the main phasealloy.

In the two-alloy blending method, Pr is preferably added to the mainphase alloy. When Pr is added to the boundary phase alloy containing alarge amount of R-rich phase, which per se is prone to oxidation,activity is further increased. Thus, oxidation occurs predominantlyduring pulverization involved in magnet production steps and in theresultant micro-powder, leading to requirement of strong countermeasuresfor preventing oxidation, or deterioration in magnet characteristicscaused by an increase in oxygen content thereof. Such countermeasuresrender the steps and apparatus for producing magnets complex, resultingin increased cost. In contrast, when Pr is added to the main phasealloy, Pr is predominantly incorporated into R₂Fe₁₄B phase, which per seis highly anti-corrosive, so that problematic oxidation can bemitigated. In addition, when Nd is partially substituted by Pr, ananisotropic magnetic field of R₂Fe₁₄B phase slightly increases. Thus,the micro-powder can be readily caused to be oriented during orientationin a magnetic field, thereby increasing magnetization and a degree oforientation of produced magnets.

As mentioned above, Pr is preferably added to the main phase alloy.However, in the course of partial substitution of Nd of the main phasealloy having a low TRE by Pr, α-Fe is prone to precipitate. One possiblereason is that the substitution by Pr increases the difference between atemperature of the liquidus at which formation of γ-Fe (high-temperaturephase) is initiated and the peritectic temperature at which formation ofR₂Fe₄B phase is initiated. Since α-Fe is difficult to pulverize,efficiency of pulverization in magnet production steps is deteriorated,thereby reducing productivity of magnets. If unpulverized α-Fe remainsin a pulverization apparatus, the composition of the resultantmicro-powder varies. If α-Fe remains in a magnet even after sintering,magnetic characteristics of the magnet are considerably deteriorated.

According to the SC method, molten metal can be supercooled at highcooling rate to a temperature lower than peritectic temperature at whichR₂Fe₁₄B phase is formed, thereby preventing precipitation of α-Fe.However, when an Nd—Fe—B ternary main phase alloy has an Nd content ofabout 28.5% by mass or less, sufficient supercooling cannot be attained,whereby α-Fe is formed. In addition, when Nd is partially substituted byPr, precipitation of α-Fe is further promoted. Thus, in order to preventprecipitation of α-Fe, the TRE of the main phase alloy must beincreased. In the two-alloy blending method, the TRE of the main phasealloy is preferably adjusted to as low a level as possible so as toenhance the mixing ratio of the boundary phase alloy.

An increase in B content is known to be remarkably effective forpreventing precipitation of α-Fe. However, when the B content of themain phase alloy increases, the B content of the boundary phase alloymust be lowered in order to adjust the total B level of the finallyproduced magnet. Addition of Co or a heavy rare earth element to themain phase alloy is also effective for preventing precipitation of α-Fe.However, when the above compositional control approaches are employed,the degree of freedom in compositional design for magnet alloydecreases. Even when the two-alloy blending method is employed, anoptimum combination of the compositions is difficult to attain.

The element Co, which exerts excellent effect for improving corrosionresistance, is preferably added to the boundary phase alloy (Kusunoki etal., T. IEEE Japan, Vol. 113-A, No. 12, 1993, 849-853). A heavy rareearth element is also confirmed to exert excellent effect for enhancingcoercivity when the element is added to the boundary phase alloy (Ito etal., Journal of the Japan Institute of the Metals, Vol. 59, No. 1(1995), 103-107).

The problems of the second aspect of the present invention are asfollows.

A series of studies were carried out on the relationship between themicrostructure of the cast R-T-B alloy ingot and the behavior uponhydrogen decrepitation or micro-pulverization, and has found thatcontrol of the dispersion state of R-rich phase is more critical, forproviding a sintered magnet alloy powder of uniform particle size, thancontrol of the crystal grain size of the alloy ingot. The inventor hasalso found that a region in which dispersion state of R-rich phase isexcessively minute (fine R-rich phase region) formed on the mold side ofthe alloy ingot is a more critical factor for controlling the particlesize of magnet powder than adverse effects of chill crystals, which arein fact contained in an alloy ingot in amounts of some % or less. Inother words, the inventor has confirmed that the percent volume of fineR-rich phase region may be in excess of 50% even when the number ofchill crystals contained in the R-T-B alloy ingot is decreased throughmodification of the composition of the alloy ingot or productionconditions; that the fine R-rich phase region deforms the particle sizedistribution of the magnet alloy powder; and that the fine R-rich phaseregion must be reduced in order to enhance magnet characteristics.

The problems to be solved by third aspect is described below.

Through the method disclosed in Japanese Patent Application No.2001-383989, reduction of fine R-rich phase region and yielding ofuniform microstructure can be attained to some extent. However, otherthan surface conditions of a roller for casting, there are a variety offactors which determine the microstructure, and such factors aredifficult to completely control during actual R-T-B alloy production.Thus, fine R-rich phase region may be formed at a portion of the alloy.

The problems to be solved by the fourth aspect are as follows.

The present inventor has carried out extensive studies on therelationship between the microstructure of the cast R-T-B alloy ingotand the behavior upon hydrogen decrepitation or micro-pulverization, andhas found that control of the dispersion state of R-rich phase is morecritical, for providing a sintered magnet alloy powder of uniformparticle size, than control of the crystal grain size of the alloyingot. The inventor has also found that a region in which dispersionstate of R-rich phase is excessively minute (fine R-rich phase region)formed on the mold side of the alloy ingot is a more critical factor forcontrolling the particle size of magnet powder than adverse effects ofchill crystals, which are in fact contained in an alloy ingot in amountsof some % or less. In other words, the inventor has confirmed that thepercent volume of fine R-rich phase region may be in excess of 50% evenwhen the number of chill crystals contained in the R-T-B alloy ingot isdecreased through modification of the composition of the alloy ingot orproduction conditions; that the fine R-rich phase region deforms theparticle size distribution of the magnet alloy powder; and that the fineR-rich phase region must be reduced in order to enhance magnetcharacteristics.

The present invention has been accomplished on the basis of this findingand an object thereof is to provide a method for producing anrare-earth-containing alloy flake, the method more effectivelypreventing formation of fine R-rich phase region in a castrare-earth-containing alloy ingot made of an R-T-B alloy, and arare-earth-containing alloy flake having a structure with excellentuniformity produced by the above method.

DISCLOSURE OF THE INVENTION

In view of the foregoing, an object of the first aspect of the presentinvention is to provide a main phase alloy for a rare earth magnet, thealloy being formed of an R-T-B alloy and to be subjected to thetwo-alloy blending method, wherein anisotropic magnetic field isenhanced and the amount of α-Fe formed is lowered at advantageously lowcost by partially substituting Nd by Pr without increasing the TRE ofthe main phase alloy for preventing precipitation of α-Fe and withoutperforming compositional control through addition of elements such as Band Co.

Accordingly, the first aspect of the present invention is directed tothe following:

-   (1) a main phase alloy for a rare earth magnet to be processed    through the two-alloy blending method, the alloy containing R(R    represents at least one rare earth element including Y) in an amount    of 26 to 30% by mass and B in an amount of 0.9 to 1.1% by mass, with    the balance being T (T represents transition metals including Fe as    an essential element), characterized in that R has a Pr content of    at least 5% by mass and the main phase alloy has a percent volume of    region containing α-Fe on the basis of the entire microstructure of    5% or less;-   (2) a main phase alloy for a rare earth magnet as described in (1),    wherein R has a Pr content of at least 15% by mass;-   (3) a main phase alloy for a rare earth magnet as described in (2),    wherein R has a Pr content of at least 30% by mass;-   (4) a main phase alloy for a rare earth magnet as described in any    one of (1) to (3), wherein at least one surface thereof has a    surface roughness, as represented by 10-point average roughness    (Rz), falling within a range of 5 μm to 50 μm;-   (5) a main phase alloy for a rare earth magnet as described in (4),    wherein at least one surface thereof has a surface roughness, as    represented by 10-point average roughness (Rz), falling within a    range of 7 μm to 25 μm;-   (6) a method for producing a main phase alloy for a rare earth    magnet as recited in any one of (1) to (5), wherein the method    comprises strip casting;-   (7) a method for producing a main phase alloy for a rare earth    magnet as described in (6), wherein the surface roughness, as    represented by 10-point average roughness (Rz), of the cast surface    of a rotating roller for casting is adjusted to fall within a range    of 5 μm to 100 μm;-   (8) a method for producing a main phase alloy for a rare earth    magnet as described in (6), wherein the surface roughness, as    represented by 10-point average roughness (Rz), of the cast surface    of a rotating roller for casting is adjusted to fall within a range    of 10 μm to 50 μm;-   (9) a method for producing a main phase alloy for a rare earth    magnet as described in any one of (1) to (3), characterized by    comprising a centrifugal casting method including depositing and    solidifying a molten metal on an inner surface of a cylindrical mold    which is rotating;-   (10) a mixed powder for a rare earth sintered magnet produced by    mixing a main phase alloy for a rare earth magnet as recited in any    one of (1) to (3) with a boundary phase alloy which has an R content    higher than that of the main phase alloy and has a Pr content of R    lower than that of the main phase alloy;-   (11) a mixed powder for a rare earth sintered magnet as described in    (10), wherein the boundary phase alloy contains substantially no Pr;    and-   (12) a rare earth sintered magnet produced through a powder    metallurgical method making use of a mixed powder for a rare earth    magnet as described in (10) or (11).

An object of the second aspect of the present invention is to provide arare earth magnet in which R-rich phase is uniformly dispersed and whichexhibits excellent magnet characteristics by suppressing formation offine R-rich phase region contained in the cast R-T-B alloy ingot, tothereby produce an alloy ingot having a microstructure of highuniformity.

Comparisons were made in terms of percent volume of fine R-rich phaseregion formed in an R-T-B alloy flake under modification of castingconditions of the SC method, particularly surface conditions of arotating roller for casting, and has found a relationship between thesurface roughness of the mold side surface of the alloy flake and thepercent volume of the formed fine R-rich phase region. The presentinvention has been accomplished on the basis of this finding.

Accordingly, the second aspect of the present invention provides thefollowing:

-   (13) a rare earth magnet alloy flake comprising an R-T-B alloy (R    represents at least one rare earth element including Y; T represents    transition metals including Fe as an essential element; and B    represents boron), characterized in that the flake has a thickness    falling within a range of 0.1 mm to 0.5 mm, and that at least one    surface of the flake has a surface roughness, as represented by    10-point average roughness (Rz), falling within a range of 5 μm to    50 μm;-   (14) a rare earth magnet alloy flake as described in (13), wherein    at least one surface of the alloy flake has a surface roughness, as    represented by 10-point average roughness (Rz), falling within a    range of 7 μm to 25 μm;-   (15) a rare earth magnet alloy flake as described in (13) or (14),    wherein the flake has a percent volume of fine R-rich phase region    in alloy that constitutes the alloy flake of 20% or less;-   (16) a method for producing a rare earth magnet alloy flake formed    of an R-T-B alloy including a strip casting method, characterized by    comprising adjusting the surface roughness, as represented by    10-point average roughness (Rz), of the cast surface of a rotating    roller for casting to fall within a range of 5 μm to 100 μm;-   (17) a method for producing a rare earth magnet alloy flake as    recited in any one of (13) to (15) formed of an R-T-B alloy    including a strip casting method, characterized by comprising    adjusting the surface roughness, as represented by 10-point average    roughness (Rz), of the cast surface of a rotating roller to fall    within a range of 5 μm to 100 μm;-   (18) a method for producing a rare earth magnet alloy flake as    described in (16) or (17), wherein the surface roughness, as    represented by 10-point average roughness (Rz), of the cast surface    of the rotating roller for casting is adjusted to fall within a    range of 10 μm to 50 μm;-   (19) a rare earth sintered magnet alloy powder produced by    subjecting, to a hydrogen descepitation step, a rare earth magnet    alloy flake as recited in any one of (13) to (15), followed by    pulverization by means of jet milling;-   (20) a rare earth sintered magnet produced from a rare earth magnet    alloy powder as recited in (19) through a powder metallurgy method;-   (21) A bonded magnet alloy powder produced by use of a rare earth    magnet alloy flake as recited in any one of (13) to (15) through an    HDDR method; and-   (22) A bonded magnet produced by use of a bonded magnet alloy powder    as recited in (21).

An object of the third aspect of the present invention is to provide amethod for producing an alloy ingot having a microstructure of highuniformity, the method more effectively preventing formation of fineR-rich phase region in the cast R-T-B alloy ingot as compared withconventional methods. Another object of the invention is to provide arare earth magnet of excellent magnet characteristics which are attainedby further increasing uniformity in the R-rich phase distribution statein the magnet.

The object of the third aspect of the present invention is as follows.

The present inventor has performed comparison in terms of percent volumeof fine R-rich phase region formed in an R-T-B alloy flake undermodification of surface conditions of a rotating roller for castingemployed in the SC method, and has found that the percent volume of theformed fine R-rich phase region depends on the morphology ofraised/dented segments provided on the mold side surface of the alloyflake, as well as on the surface roughness of the mold side surface ofthe alloy flake. The present invention has been accomplished on thebasis of this finding.

Accordingly, the third aspect of the present invention provides thefollowing:

-   (23) a rare-earth-containing alloy flake, characterized in that the    alloy flake has a thickness falling within a range of 0.1 mm to 0.5    mm; at least one surface of the alloy flake has a plurality of    elongated raised/dented segments (i.e., small ridge/valley areas)    formed so as to cross one another; and the surface having the    elongated raised/dented segments has a surface roughness, as    represented by 10-point average roughness (Rz), falling within a    range of 3 μm to 30 μm;-   (24) a rare-earth-containing alloy flake as described in (23),    wherein the alloy flake comprises an R-T-B alloy (R represents at    least one rare earth element including Y; T represents transition    metals including Fe as an essential element; and B represents boron)    which serves as a raw material for producing a rare earth magnet;-   (25) a rare-earth-containing alloy flake as described in (24),    wherein the flake has a percent volume of fine R-rich phase region    in alloy that constitutes the alloy flake of 20% or less;-   (26) a method for producing a rare-earth-containing alloy flake    including a strip casting (SC) method, characterized by comprising    employing a rotating roller for casting, the roller having, on the    cast surface, a plurality of elongated raised/dented segments formed    so as to cross one another and having a surface roughness of the    cast surface, as represented by 10-point average roughness (Rz),    falling within a range of 3 μm to 30 μm;-   (27) a method for producing a rare-earth-containing alloy flake as    described in (26), wherein the alloy flake has a thickness falling    within a range of 0.1 mm to 0.5 mm; at least one surface of the    alloy flake has a plurality of elongated raised/dented segments    formed so as to cross one another; and the surface having the    elongated raised/dented segments has a surface roughness, as    represented by 10-point average roughness (Rz), falling within a    range of 3 μm to 30 μm;-   (28) a method for producing a rare-earth-containing alloy flake as    described in (26) or (27), wherein the rare-earth-containing alloy    flake comprises an R-T-B alloy (R represents at least one rare earth    element including Y; T represents transition metals including Fe as    an essential element; and B represents boron) which serves as a raw    material for producing a rare earth magnet;-   (29) a method for producing a rare-earth-containing alloy flake as    described in (28), wherein the flake has a percent volume of fine    R-rich phase region in alloy that constitutes the alloy flake of 20%    or less;-   (30) an alloy powder for a rare earth sintered magnet produced by    subjecting, to a hydrogen decrepitation step, a    rare-earth-containing alloy flake as recited in (24) or (25),    followed by pulverization by means of a jet mill;-   (31) a rare earth sintered magnet produced from an alloy powder for    a rare earth sintered magnet as recited in (30) through a powder    metallurgy method;-   (32) an alloy powder for a bonded magnet produced by use of a    rare-earth-containing alloy flake as recited in (24) or (25) through    an HDDR method; and-   (33) a bonded magnet produced by use of an alloy powder for a bonded    magnet as recited in (32).

The object of the fourth aspect of the present invention is summarizedas follows.

The present inventor has previously carried out extensive studies on therelationship between the microstructure of the cast R-T-B alloy ingotand the behavior upon hydrogen decrepitation or micro-pulverization, andhas found that control of the dispersion state of R-rich phase is morecritical, for providing a sintered magnet alloy powder of uniformparticle size, than control of the crystal grain size of the alloyingot. The present inventor has performed comparison in terms of percentvolume of fine R-rich phase region formed in an R-T-B alloy flake undermodification of casting conditions of the SC method, particularlysurface conditions of a rotating roller for casting, and has found arelationship between the surface roughness of the mold side surface ofthe alloy flake and the percent volume of the formed fine R-rich phaseregion. The present inventor has accomplished a method for producingalloy flakes having a microstructure of high uniformity and a percentageof fine R-rich phase region of 20% or less.

Also the present inventor has found that uniformity of the structure canbe improved more effectively by providing a plurality of elongatedraised/dented segments constituting surface roughness of a rotatingroller for casting so as to cross one another.

That is, the present inventor has found that reduction in fine R-richphase region and improvement in uniformity of the structure can beachieved by the above method and that uniformity of the structure can beimproved more effectively by controlling not only the surface roughnessexpressed by a numerical value of the rotating roller, but also themorphology of raised/dented segments as an origin of the surfaceroughness.

The present invention has been accomplished on the basis of the findingobtained by previous studies and the following extensive studies and anobject thereof is to provide a method for producing anrare-earth-containing alloy flake, the method more effectivelypreventing formation of fine R-rich phase region in a castrare-earth-containing alloy ingot made of an R-T-B alloy by furtherimproving the morphology of raised/dented segments on the rotatingroller for casting, and a rare earth sintered magnet of excellent magnetcharacteristics which are attained by further increasing uniformity inthe R-rich phase distribution state in the magnet.

Accordingly, the fourth aspect of the present invention provides thefollowing:

-   (34) a method for producing a rare-earth-containing alloy flake    including a strip casting method, characterized by comprising    employing a rotating roller for casting, the roller having, on the    cast surface, a plurality of elongated raised/dented segments and    having a surface roughness provided by a plurality of elongated    raised/dented segments, as represented by 10-point average roughness    (Rz), falling within a range of 3 μm to 60 μm, 30% or more of    raised/dented segments among entire elongated raised/dented segments    extending in a direction forming an angle of 30° or more to a roller    rotation direction;-   (35) a method for producing a rare-earth-containing alloy flake as    described in (34), characterized by comprising employing a rotating    roller for casting, 30% or more of raised/dented segments among    entire elongated raised/dented segments extending in a direction    forming an angle of 45° or more to a roller rotation direction;-   (36) a method for producing a rare-earth-containing alloy flake as    described in (34), characterized by comprising employing a rotating    roller for casting, 50% or more of raised/dented segments among    entire elongated raised/dented segments extending in a direction    forming an angle of 30° or more to a roller rotation direction;-   (37) a method for producing a rare-earth-containing alloy flake as    described in any one of (34) to (36), characterized by comprising    employing a rotating roller for casting, 50% or more of    raised/dented segments among entire elongated raised/dented segments    extending in a direction forming an angle of 45° or more to a roller    rotation direction;-   (38) a method for producing a rare-earth-containing alloy flake as    described in any one of (34) to (37), wherein the    rare-earth-containing alloy flake comprising an R-T-B alloy (R    represents at least one rare earth element including Y; T represents    transition metals including Fe as an essential element; and B    represents boron) which serves as a raw material for producing a    rare earth magnet in the production of the rare-earth-containing    alloy flake by the strip casting method;-   (39) an alloy flake for rare earth magnet produced by the method as    recited in (38), which has a percent volume of fine R-rich phase    region in alloy that constitutes the alloy flake of 20% or less;-   (40) a powder for a rare earth sintered magnet produced by    subjecting, to a hydrogen decrepitation step, an alloy flake for    rare earth magnet produced by the method as recited in (38),    followed by pulverization by means of a jet mill;-   (41) a rare earth sintered magnet produced from an alloy powder for    a rare earth sintered magnet as recited in (7) through a powder    metallurgy method;-   (42) an alloy powder for a bonded magnet produced by use of an alloy    flake for rare earth magnet produced by the method as recited    in (38) through an HDDR method; and-   (43) a bonded magnet produced by use of an alloy powder for a bonded    magnet as recited in (42).

BRIEF DESCRIPTION OF DRAWINGS

FIG. 1 is a photograph showing a cross section of an alloy flake for arare earth magnet produced through a conventional SC method.

FIG. 2 is a photograph showing a cross section of an alloy flake for arare earth magnet produced through the SC method according to thepresent invention.

FIG. 3 is the photograph of FIG. 1, but α-Fe-containing region isenclosed by the dotted line.

FIG. 4 shows photographs showing cross sections of an alloy flake for arare earth magnet produced through the centrifugal casting methodaccording to the present invention.

FIG. 5 is a sketch of a casting apparatus employed in a strip casting(SC) method.

FIG. 6 is a sketch of a centrifugal casting apparatus for sprinklingmolten alloy by centrifugal force and depositing the sprinkled alloy onthe inner wall of the mold.

FIG. 7 shows the microstructure of a cross section of a rare earthmagnet alloy flake containing fine R-rich phase region produced througha conventional SC method.

FIG. 8 shows the microstructure of a cross section of a rare earthmagnet alloy flake according to the third Aspect of the presentinvention.

FIG. 9 shows the microstructure of the same observation area as that ofFIG. 1, but the boundary between the fine R-rich phase region and thenormal portion is shown by the dotted line.

FIG. 10 shows the microstructure of a cross section of a rare earthmagnet alloy flake according to the Fourth Aspect of the presentinvention.

BEST MODE FOR CARRYING OUT THE INVENTION

Hereinafter, the present invention is described in the order from thefirst aspect to the fourth aspect.

First Aspect

As described above, when a rare earth magnet formed of R-T-B alloy isproduced through the two-alloy blending method, Pr is preferably addedto the main phase alloy for incorporating Pr into R. However, when Nd ispartially substituted by Pr, precipitation of α-Fe is promoted.Therefore, in conventional techniques, another element must be added tothe main phase alloy, or the composition of the main phase alloy must bemodified. The above compositional control approaches greatly limit thedegree of freedom in designing of alloy, which is a merit of thetwo-alloy blending method. Thus, effective use of Pr in production ofrare earth magnets through the two-alloy blending method cannot be fullyattained.

The present inventor has found that precipitation of α-Fe can be greatlysuppressed by modifying casting conditions for casting a main phasealloy for a rare earth magnet through the SC method, particularly bymodifying surface conditions of a rotating roller for casting, therebyimproving heat transfer from molten alloy to the roller. Thus, theinventor has successfully attained, in the two-alloy blending method,substitution of a large number of Nd atoms contained in a low-TRE mainphase alloy by Pr.

The present inventor has confirmed that precipitation of α-Fe in a mainphase alloy for a rare earth magnet produced through a centrifugalcasting method including depositing and solidifying a molten metal on aninner surface of a cylindrical mold which is rotating is moreeffectively inhibited, as compared with a conventional SC method, bymeans of lowering the deposition rate. The inventor has successfullyproduced, through the centrifugal casting method, a main phase alloy foruse in the two-alloy blending method having a small TRE and a high Prcontent, with formation of α-Fe being greatly suppressed. The inventorhas also found that a casting method including pouring a molten alloyinto a rotary body, sprinkling the molten alloy by rotating the rotarybody, and depositing and solidifying the sprinkled molten alloy on aninner surface of the cylindrical mold which is rotating (hereinafterreferred to as a centrifugal force-sprinkled molten alloy-typecentrifugal casting method) is particularly preferred as the centrifugalcasting method.

The present inventor has also confirmed that, when no difference inamount of Pr added in a rare earth magnet (final product) exists betweenproducts, oxidation occurring during magnet production steps isinhibited more effectively for a rare earth magnet produced through thetwo-alloy blending method employing a main phase alloy according to thepresent invention containing most portions of Pr than for a rare earthmagnet produced through a customary single-alloy method, therebylowering processing cost, lowering the oxygen content of rare earthmagnets, and improving characteristics of the magnets.

FIG. 1 shows a typical back-scattered electron image, observed under anSEM (scanning electron microscope), of a cross-section of a main phasealloy (TRE: 28.5% by mass, Nd:Pr=1:1 (in R), B: 1% by mass; and balance:Fe) flake for a rare earth magnet which has been cast through aconventional SC method.

In FIG. 1, the left side corresponds to the mold side, and the rightside to the free surface side. On the mold side, the alloy flake has asurface roughness, as represented by 10-point average roughness (Rz), of3.4 μm. In FIG. 1, white areas correspond to R-rich phase which assumesthe form of lamellar portions extending in the thickness direction, orthe form of a small pool of oriented lamellar fragments.

During casting of a main phase alloy through the SC method, a rare earthcomponent contained in the molten alloy is oxidized to form acorresponding oxide, as the oxygen content of the atmosphere increases.The resultant oxide serves as a nucleation site for forming α-Fe,whereby precipitation of α-Fe is promoted. In another case, when thetemperature of the molten alloy is lowered to a level near the liquidus,α-Fe is formed in the molten alloy. In still another case, when the castflake is thick, the solidification rate particularly on the free surfaceside decreases, thereby readily forming α-Fe. Thus, in the above SCmethod, the oxygen content and temperature of the molten alloy arerigorously controlled during casting, and the thickness of cast flakesis controlled to as slightly thin as about 0.2 mm, so as to preventformation of α-Fe. However, as shown in FIG. 1, α-Fe is precipitated insome sites on the free surface side of a main phase alloy flake for arare earth magnet cast through the conventional SC method. In theback-scattered electron image shown in FIG. 1, α-Fe is observed as aportion of higher tint than that of R₂Fe₁₄B phase (main phase),specifically as black dots in FIG. 1.

FIG. 2 shows a back-scattered electron image of a flake of an alloyhaving a composition identical to that of the above alloy (FIG. 1), theflake having been produced through the improved SC method according tothe present invention. In FIG. 2, the left side corresponds to the moldside, and the right side to the free surface side. The SC methodemployed in the present invention is characterized in that formation ofα-Fe in an alloy is prevented by controlling the surface roughness ofthe mold side of a flake produced from the alloy. As shown in FIG. 2,the alloy flake produced through the improved SC method of the presentinvention contains no α-Fe, and homogeneity in R-rich phase dispersionstate is satisfactory from the mold side to the free surface side.

Even when a conventional SC method is employed, the produced alloyflakes include, to some extent, those having a uniform microstructurefree from α-Fe as shown in FIG. 2. However, alloy flakes containing α-Feas shown in FIG. 1 are also produced simultaneously. Thus, the percentvolume of α-Fe-containing region in the entire microstructure of themain phase alloy cannot be decreased to 5% or less. Difference inportions of microstructure of the alloy produced through a conventionalSC method may be attributable to difference in conditions of contactbetween the roller surface and the molten alloy or alloy flake; e.g.,the fine surface state of the rotating roller for casting, molten alloysupply conditions, and the atmosphere during casting.

The percent volume of α-Fe-containing region in the microstructure of amain phase alloy can be determined in the following manner. FIG. 3 is aback-scattered electron image of the same observation area as that ofFIG. 1, but in FIG. 3 the α-Fe-containing region is enclosed by theline. Since α-Fe is precipitated as grains or dendrites over acomparatively wide area extending some tens of μm or more and theboundary between regions can be readily identified, the percent area ofthe α-Fe-containing region in the observation area can be calculated bymeans of a graphic image analyzer. The percent area in the cross-sectioncorresponds to the percent volume of the alloy. As mentioned above, thepercent volume of α-Fe-containing region varies depending on conditionsof contact between the surface of a rotating roller for casting and themolten alloy or alloy flake as well as on the composition of the alloy.In addition, the surface of a rotating roller for casting is notcompletely uniform, and even a subtle change in amount of poured moltenalloy varies the conditions of contact between the rotating roller forcasting and the molten alloy. Therefore, the percent volume ofα-Fe-containing region greatly varies among alloy flakes or within onealloy flake, even when the alloy flakes are produced under the sameconditions. Thus, graphic image analysis is performed by use of about 5to about 10 flakes under a wide observation area at a low magnificationof about 100 to about 200 times, and obtained percent area values areaveraged, to thereby calculate the percent volume of the α-Fe-containingregion for the entirety of the alloy.

The relationship between the effect for preventing precipitation of α-Feand the surface roughness of the mold side surface of an alloy flakeproduced through the SC method can be described as follows. In order toobtain a smooth mold side surface of an alloy flake, the surface of arotating roller for casting must be smooth and have high wettabilitywith respect to the molten alloy. When such a rotating roller isemployed, heat is transferred from the molten alloy to the mold atremarkably high efficiency (i.e., heat transfer coefficient is high).Thus, the mold side alloy is rapidly cooled, solidified, and reduced insize, allowing rising, or exfoliation, of portions of the alloy from thesurface of the rotating roller for casting. The heat transfer from theresultant portion to the roller greatly decreases, whereby the coolingefficiency thereafter is greatly decreases. Such a large decrease insolidification rate is considered to cause precipitation of α-Fe on thefree surface side. Such a phenomenon is prone to occur, when the surfaceroughness of the rotating roller for casting is less than 5 μm.

According to the method for producing a main phase alloy for a rareearth magnet of the present invention including an SC method, thesurface roughness, as represented by 10-point average roughness (Rz), ofthe surface of a rotating roller for casting is controlled to fallwithin a range of 5 μm to 100 μm, preferably within a range of 10 μm to50 μm. When the surface of the rotating roller for casting is controlledto fall within the above range, the minute irregularities formed on thesurface of the rotating roller for casting cannot be filled completelywith the molten alloy, because of its viscosity. Thus, many portions ofthe alloy remain not in contact with the roller, thereby lowering theheat transfer coefficient. Although an excessive decrease in heattransfer coefficient induces precipitation of α-Fe, a surface roughnessfalling within the above range prevents precipitation of α-Fe, therebyappropriately controlling the heat transfer coefficient.

Through provision, on the surface of a rotating roller for casting, of asurface roughness falling within the above range, excessive heattransfer at an initial stage of solidification of molten alloy can beprevented, thereby preventing reduction in size of flakes formed of thealloy caused by rapid solidification. In addition, engagement of thesurface of the rotating roller for casting to the irregularities of thealloy flake is also effective for preventing partial peeling of thealloy flake from the roller caused by solidification/reduction in size.Thus, formation of α-Fe in the alloy is considered to be preventedthrough a decreased change in solidification rate in a range of the moldside where solidification of the molten alloy begins to the free surfaceside where solidification complete.

When the surface roughness of the rotating roller for casting increases,the surface roughness of the mold side surface of the alloy flakenecessarily increases, through transfer of the irregularities of theroller to the mold side surface of the alloy flake to some extent.According to the present invention, the surface roughness, asrepresented by 10-point average roughness (Rz), of the surface of arotating roller for casting is controlled to fall within a range of 5 μmto 100 μm, preferably within a range of 10 μm to 50 μm. Thus, at leastone surface of the as-cast alloy product has a surface roughness fallingwithin a range of 5 μm to 50 μm, preferably 7 μm to 25 μm.

However, when the surface roughness of the rotating roller for castingis in excess of 100 μm, the surface roughness can be filled with themolten alloy, thereby increasing heat transfer coefficient and furtherincreasing the surface roughness of the mold side of the produced alloyflake. In this case, although exfoliation of the alloy flake can besuppressed, uniformity in dispersion of R-rich phase cannot be attaineddue to large surface roughness, which is disadvantageous.

Another mode of the method for producing a main phase alloy for a rareearth magnet according to the present invention is a centrifugal castingmethod. For example, a molten alloy is sprinkled by a rotary body suchas a rotatable tundish, to thereby greatly decreasing the rate forsupplying a molten alloy. Thus, even when the alloy has a low TRE, amain phase is epitaxially grown, thereby preventing formation of α-Fe.FIG. 4 is a back-scattered electron image of a cross-section of an alloyhaving a composition identical to the alloy shown in FIG. 1 or 2 andproduced through the above centrifugal casting method. In FIG. 4, theupper photograph shows a microstructure of a cross-section at a levelfrom the mold side of 0.5 mm, the middle photograph shows amicrostructure of a cross-section at a middle level, and the lowerphotograph shows a microstructure of a cross-section at a level from thefree surface side of 0.5 mm. As is clear from FIG. 4, the main phasealloy having a composition identical to the alloy shown in FIG. 1 or 2and produced through the above centrifugal casting method has amicrostructure in which no α-Fe from the mold side to the free surfaceside is contained and R-rich phase is dispersed remarkably uniformly.

The elements of the present invention will next be described in detail.

(11) TRE in a Main Phase Alloy for a Rare Earth Magnet

According to the present invention, the TRE in the main phase alloy fora rare earth magnet is controlled to fall within a range of 26 to 30% bymass. When a sintered magnet is produced through the two-alloy blendingmethod, the main phase alloy employed in the two-alloy blending methodpreferably has a low TRE for increasing the ratio of the boundary phasealloy in the mixture and facilitate mixing of the main phase alloy withthe grain boundary phase. In consideration of the fact thathigh-performance magnet produced by use of R-T-B alloy generally has aTRE of about 32% by mass or less, the main phase alloy should have a TREof 30% by mass or less, preferably 29% by mass or less. When the alloyhas a stoichiometry of Nd₂Fe₁₄B, Nd content accounts for 26.7% by mass.When the TRE is less than the level, precipitation of α-Fe inevitablyoccurs. Therefore, a great decrease in TRE from the level is notpossible, and the lower limit is 26% by mass. TRE is preferably 27% bymass or more.

(12) Pr Content of R

According to the present invention, the Pr content of R contained in themain phase alloy for a rare earth magnet is controlled to at least 5%.Conventionally, a main phase alloy having a small TRE to be processed bythe two-alloy blending method suffers from problematic precipitation ofα-Fe, when Nd is partially substituted by Pr (in R). Alternatively, Pris added to the main phase alloy only when formation of α-Fe has beensuppressed by controlling the levels of other components; i.e., byincreasing the B content or adding Co. However, according to theimproved SC method of the present invention or a centrifugal castingmethod which allows feeding of a molten alloy at such a low speed thatthe molten alloy is sprinkled by centrifugal force, precipitation ofα-Fe is effectively suppressed. Even when Pr is added, α-Fe is difficultto precipitate. Thus, the Pr content of R can be controlled to at least5%, even though the above component adjustment is omitted. In addition,the main phase alloy for a rare earth magnet of the present inventionhas a large degree of freedom in compositional design for magnet alloy,since the levels of other components is has not been modified; i.e.,modification of B content or addition of Co has not been performed. Inorder to enhance the effect of addition of Pr, which is onecharacteristic of the present invention, the Pr content of R iscontrolled to preferably at least 15% by mass, more preferably at least30% by mass.

(13) Percent Volume of α-Fe-Containing Region

According to the present invention, the main phase alloy has a percentvolume of region containing α-Fe on the basis of the entiremicrostructure of 5% or less. α-Fe deteriorates efficiency ofpulverizing the main phase alloy and induces variation in composition ofthe alloy. If α-Fe remains in magnets, magnetism is deteriorated. Whenthe percent volume of the α-Fe-containing region is in excess of 5%,these drawbacks become critical. Through production by theaforementioned SC method or centrifugal casting method according to theinvention, the percent volume of α-Fe-containing region in themicrostructure of the main phase alloy of the present invention can becontrolled to 5% or less, even when the Pr content of R is controlled toat least 5% by mass, preferably 15% by mass, more preferably 30% bymass.

(14) Surface Roughness of the Mold Side of an Alloy Flake ProducedThrough SC Method

According to the present invention, the surface roughness of the mainphase alloy flake produced through the strip casting method, asrepresented by 10-point average roughness (Rz), falls within a range of5 μm to 50 μm. As mentioned above, when the surface roughness is 5 μm orless, variation in heat transfer between the roller for casting and thealloy flake increases during solidification, resulting in variation ofthe rate of solidification of the molten alloy. Thus, α-Fe isprecipitated at some portions. When the surface roughness is 50 μm ormore, although precipitation of α-Fe is inhibited, uniformity in R-richphase dispersion is failed to be attained, which is disadvantageous.

Thus, the surface roughness of the mold side of the alloy flakepreferably falls within a range of 7 μm to 25 μm.

Herein, the term “surface roughness” refers to a surface roughnessdetermined under the conditions specified in JIS B 0601 “Surfaceroughness—Definitions and Designation,” and 10-point average roughness(Rz) is defined therein. Specifically, a surface to be measured is cutwith a plane which is perpendicular thereto, to thereby obtain a contourappearing on a cut end (profile curve). Any surface waviness componentlonger than a prescribed wavelength is cut off from the profile curve bymeans of a phase-compensation-type high-pass filter or a similar device,to thereby obtain a curve (roughness curve). Only the reference lengthis sampled from the roughness curve in the direction of its mean line,and the sum of the average value of absolute values of the heights ofthe five highest profile peaks (Yp) and the depths of the five deepestprofile valleys (Yv) measured in the vertical direction from the meanline of this sampled portion is calculated, to thereby obtain the10-point average roughness (Rz). Measurement parameters such asreference length are defined in the above JIS B 0601, as standard valuesof reference length for determining corresponding surface roughnessvalues.

Since the surface roughness of the mold side of an alloy flake oftenvaries in a wide range among samples to be measured, an average value ofsurface roughness for at least five flakes should be employed.

(15) Strip Casting (SC) Method

The method of producing R-T-B alloy for a rare earth magnet through thestrip casting method will be described with reference to a sketch of anapparatus shown in FIG. 5.

Generally, a rare earth alloy is made molten by use of a refractorycrucible 1 in vacuum or an inert gas atmosphere, because it is highlyactive. The thus-molten alloy is maintained at 1,350 to 1,500° C. for apredetermined period of time, and supplied, via a tundish 2 havingoptional flow-control means or slag-removing means, to a rotating roller3 for casting whose interior is cooled with water. The rate of supplyingthe molten alloy and the rotation speed of the rotating roller areappropriately regulated in accordance with the thickness of the alloyflakes to be produced. Generally, the rotation speed of the rotatingroller is about 1 to about 3 m/s (in terms of peripheral velocity). Therotating roller for casting is preferably made of copper or copperalloy, from the viewpoint of high thermal conductivity and availability.The surface of the rotating roller for casting is prone to adsorbmetallic material, depending on the type of alloy to be produced andsurface conditions of the rotating roller. Thus, provision of anoptional cleaning apparatus stabilizes qualities of the cast alloy.

The alloy 4 solidified on the rotating roller is released from theroller on the side opposite the tundish side and collected into acollection container 5. The conditions of R-rich phase can be controlledby means of heating/cooling means provided in the collection container(Japanese Patent Application Laid-Open (kokai) Nos. 09-170055 and10-36949).

(16) Centrifugal Casting Method

Similar to a conventional centrifugal casting method, the centrifugalcasting method employed in the present invention includes feeding amolten metal into the interior of a cylindrical mold which is rotating,to thereby simultaneously deposit and solidify the molten metal.However, through employment of a conventional method in which the moltenmetal is caused to fall simply by gravity from holes provided in thetunduish, the rate of depositing molten metal is difficult to decrease,possibly precipitating α-Fe in the alloy. Thus, a casting methodsuitable for the present invention include feeding a molten alloy into arotary body, forming droplets of the molten alloy by application ofcentrifugal force, and sprinkling the droplets, thereby depositing thealloy on the inner wall of the mold. Such a method can greatly decreasethe deposition rate and increase solidification rate, and exerts greatereffect for preventing precipitation of α-Fe, as compared with aconventional SC method (Japanese Patent Application No. 2000-262605).

FIG. 6 is a sketch of an apparatus employed in a centrifugal castingmethod for sprinkling molten alloy by centrifugal force and depositingthe alloy on the inner wall of the mold. Generally, a rare earth alloyis made molten by use of a refractory crucible 6 in vacuum or an inertgas atmosphere, because it is highly active. The thus-molten alloy ismaintained at 1,350 to 1,500° C. for a predetermined period of time, andsupplied, via a trough 7, to a rotary body 8. The molten alloy is causedto be sprinkled to the inner wall of a cylindrical mold 9 throughrotation of the rotary body 8. Thus, the rate for pouring the moltenalloy can be controlled, to thereby produce an alloy 10 at a desireddeposition rate. If the axis of rotation of the rotary body 8 and theaxis of rotation of the mold 9 form a certain angle, deposition area canbe expanded over the inner wall of the mold in a longitudinal direction,thereby controlling the deposition rate of the molten metal.

(17) TRE and Pr Content of Boundary Phase Alloy

The main phase alloy for a rare earth magnet produced according to thepresent invention and to be subjected to the two-alloy blending methodis mixed with a separately produced boundary phase alloy for a rareearth magnet to be subjected to the two-alloy blending method. Theresultant mixture is pulverized, molded, and sintered, to therebyproduce high-performance anisotropic magnets.

The boundary phase alloy predominantly provides R-rich grain boundaryphase rather than R₂T₁₄B phase (main phase), and therefore, has a TREhigher than that of the main phase alloy. According to the presentinvention, the Pr content of R contained in the boundary phase alloy islower than that in the main phase alloy. As mentioned above, Pr ispreferably added in a larger amount to the main phase alloy and in anamount as small as possible to the boundary phase alloy, from aviewpoint of corrosion resistance and orientation in a magnetic field.More preferably, the boundary phase alloy contains no substantial Pr,other than impurities present in the raw material.

(18) Methods for Producing Mixed Powder for a Rare Earth Sintered Magnetand Producing Rare Earth Magnets

The main phase alloy and the boundary phase alloy for a rare earthmagnet according to the present invention are typically performedsequentially in the order of hydrogen decrepitation andmicro-pulverization, to thereby produce alloy powder having a size ofapproximately 3 μm (FSSS). Hydrogen decrepitation includes a hydrogenabsorption step as a first step and a hydrogen desorption step as asecond step. In the hydrogen absorption step, hydrogen is caused to beabsorbed predominantly in the R-rich phase of alloy flakes in a hydrogengas atmosphere at 266 hPa to 0.3 MPa·G. The R-rich phase is expanded involume due to R hydride generated in this step, to thereby minutelybreak the alloy flakes themselves or generate numerous micro-cracks.Hydrogen absorption is carried out within a temperature range of ambienttemperature to approximately 600° C. However, in order to increaseexpansion in volume of R-rich phase so as to effectively reduce theflakes in size, hydrogen absorption is preferably performed underincreased pressure and within a temperature range of ambient temperatureto approximately 100° C. The time for hydrogen absorption is preferablyone hour or longer. The R hydride formed through the hydrogen absorptionstep is unstable in the atmosphere and readily oxidized. Thus, thehydrogen-absorbed product is preferably subjected to hydrogen desorptiontreatment by maintaining the alloy flakes at about 200 to about 600° C.in vacuum of 1.33 hPa or less. Through this treatment, R hydride can betransformed into a product stable in the atmosphere. The time forhydrogen desorption treatment is preferably 30 minutes or longer. If theatmosphere is controlled for preventing oxidation during steps to becarried out after hydrogen absorption to sintering, hydrogen desorptiontreatment can also be omitted.

Micro-pulverization is a step of pulverizing alloy flakes for attaininga particle size of approximately 3 μm (FSSS). Among pulverizers forperforming the micro-pulverization, a jet mill is most preferred,because of attaining high productivity and a sharp particle sizedistribution profile. Upon micro-pulverization, the atmosphere iscontrolled to an inert gas atmosphere such as an argon gas atmosphere ornitrogen gas atmosphere. The inert gas may contain oxygen in an amountof 2% by mass or less, preferably 1% by mass or less. The presence ofoxygen enhances pulverization efficiency and attains oxygenconcentration of the powder produced through pulverization to 1,000 to10,000 ppm, to thereby enhance oxidation resistance of the alloy powder.In addition, abnormal grain growth during sintering can also beprevented.

The main phase alloy and the boundary phase alloy are mixed at apredetermined compositional ratio. Mixing may be performed at any step;i.e., before hydrogen decrepitation, before micro-pulverization, orafter micro-pulverization. When the two alloys are considerablydifferent from each other in terms of pulverizability, mixing ispreferably performed after completion of micro-pulverization. However,when only a small difference in terms of pulverizability is foundbetween the two alloys, mixing may be performed before hydrogendecrepitation.

When the alloy powder for a rare earth magnet is molded in a magneticfield, in order to reduce friction between the powder and the inner wallof a mold and to reduce friction generated among powder particles forenhancing orientation, a lubricant such as zinc stearate is preferablyadded to the powder. The amount of the lubricant to be added is 0.01 to1% by mass. Although the lubricant may be added before or aftermicro-pulverization, the lubricant is preferably mixed sufficiently,before molding in magnetic field, in an inert gas atmosphere such asargon gas or nitrogen gas by use of a mixing apparatus such as aV-blender.

The powder having a particle size of about 3 μm (FSSS) obtained throughmicro-pulverization is press-molded in magnetic field by use of amolding apparatus. The mold to be employed is fabricated from a magneticmaterial and a non-magnetic material in combination in consideration ofthe orientation of magnetic field in the mold cavity. The pressure atmolding is preferably 0.5 to 2 t/cm², and the magnetic field in the moldcavity during molding is preferably 5 to 20 kOe. The atmosphere duringmolding is preferably an inert gas atmosphere such as argon gas ornitrogen gas. However, if the powder has been subjected to theaforementioned anti-oxidation treatment, molding can be performed inair. Molding may be performed through cold isostatic pressing (CIP) orrubber isostatic pressing (RIP) employing a rubber mold. Since the alloypowder is pressed isostatically through CIP or RIP, variation inorientation of magnetization during press-molding is lowered. Thus, thedegree of orientation of the produced compact can be increased ascompared with that produced by use of a metal mold, and maximum magneticenergy product can be enhanced.

Sintering of the compact for a rare earth magnet is performed at 1,000to 1,100° C. Prior to reaching the sintering temperature, lubricant andhydrogen contained in the micro-powder must be removed as completely aspossible. The lubricant is removed by maintaining the compact preferablyunder the conditions: in vacuum of 1.33×10⁻² hPa or under an argon gasflow atmosphere at reduced pressure; at 300 to 500° C.; and for 30minutes or longer. Hydrogen is removed by maintaining the compactpreferably under the conditions: in vacuum of 1.33×10⁻² hPa or less; at700 to 900° C.; and for 30 minutes or longer. The atmosphere duringsintering is preferably an argon gas atmosphere or a vacuum atmosphereof 1.33×10⁻² hPa or less. A retention time at the sintering temperatureof one hour or longer is preferred.

After completion of sintering, in order to enhance the coercivity ofsintered magnet to be produced, the sintered product may be treated at500 to 650° C. in accordance with needs. An argon gas atmosphere or avacuum atmosphere is preferred, and a retention time of 30 minutes orlonger is preferred.

The rare earth magnet produced through the two-alloy blending method byuse of a Pr-containing main phase alloy of the present invention is moreexcellent than a similar rare earth magnet produced through thesingle-alloy method by use of a raw material alloy containing a similaramount of Pr. The following is the conceivable advantages of the formerrare earth magnet.

When the single-alloy method is employed, the composition of the rawmaterial alloy for a rare earth magnet coincides approximately with thatof the rare earth magnet (final product), and the difference between twocompositions may be attributable to subtle compositional variationprovided during production steps. The TRE is about 31 to about 33% bymass. The alloy powder has an R-rich phase content of about 5 to about10%. When R is predominantly comprising Nd, Pr is prone to form R-richphase rather than R₂T₁₄B phase. Thus, Pr content becomes higher in theR-rich phase than in the entirety of the alloy. Therefore, R-rich phase,which per se is active is further activated, and oxidation readilyoccurs during pulverization involved in magnet production steps and inthe resultant micro-powder. The excessively increased oxygen contentdeteriorates magnet characteristics. When the Pr content increases,thorough countermeasures for preventing oxidation during steps isrequired. Such countermeasures result in increased cost and decrease inproduction efficiency. In addition, when the Pr content of R-rich phaseremains high after formation of a sintered magnet, corrosion resistanceof the magnet decreases due to active R-rich phase which may also bepresent in the alloy or micro-powder.

In contrast, in the rare earth magnet according to the present inventionproduced through the two-alloy blending method, Pr is provided from themain phase alloy. Thus, the Pr content is high in R₂T₁₄B phase whichoriginally contains Pr and is low in R-rich phase. During sintering, Prmay diffuse from R₂T₁₄B phase to R-rich phase, resulting in slightincrease in Pr content of R-rich phase. Although the Pr content ofR-rich phase slightly increases, increase in Pr content of R-rich phasecan be suppressed more effectively as compared with a magnet producedthrough the single-alloy method employing a raw material alloy having ahigh Pr content of R-rich phase, whereby corrosion resistance can beimproved. In addition, since the main phase alloy included in the rawmaterial has a high Pr content, an anisotropic magnetic field in R₂T₁₄Bphase increases and ease of orientation during orientation in a magneticfield can be improved, thereby increasing magnetization and a degree oforientation of produced magnets.

EXAMPLES Example 11

Neodymium, praseodymium, ferroboron, aluminum, and iron were mixed tothereby obtain the following alloy composition: TRE: 28.5% by mass(Nd:Pr=1:1 (in R)); B: 1.00% by mass; Al: 0.30% by mass; and a balanceof iron. The resulting mixture was melted in an alumina crucible in anargon gas atmosphere (1 atm) by use of a high-frequency inductionmelting furnace. The resulting molten alloy was cast through stripcasting, to thereby prepare alloy flakes.

The roller for casting having a diameter of 300 mm and made of purecopper was employed. During casting, the inside of the copper roller wascooled by water. The roller had a cast surface roughness, as representedby 10-point average roughness (Rz), of 20 μm and was rotated at aperipheral velocity of 0.9 m/s, to thereby produce alloy flakes having amean thickness of 0.26 mm.

The thus-produced alloy flakes were found to have a surface (mold side)roughness, as represented by 10-point average roughness (Rz), of 9 μm.Ten flakes were selected from the alloy flakes and polished in a fixedstate. Each flake was observed under a scanning electron microscope(SEM) and a back-scattered electron image (BEI) was captured at amagnification of ×200. Through analysis of the thus-captured photographby means of an image graphic analyzer, the percent volume ofα-Fe-containing region was found to be 1% or less.

Example 12

An alloy having a composition similar to that of the alloy of Example 11was melted in an alumina crucible in an argon gas atmosphere by use of ahigh-frequency induction melting furnace. The resulting molten alloy wascast by use of a centrifugal casting apparatus including a rotatabletundish.

During casting, the molten alloy was deposited on the inner wall of themold at an average deposition rate of 0.01 cm/s. The rotation rate ofthe mold was modified such that centrifugal force is adjusted to 3 G.Centrifugal force (about 20 G) was applied to the molten alloy containedin the rotatable tundish, to thereby sprinkle the molten alloy.

The thus-produced alloy flakes were found to have a thickness of 7 to 10mm. From each alloy flake, each sample cut at levels in the thicknessdirection of 7 mm, 8.5 mm, and 10 mm was polished in a fixed state. Eachflake was observed under a scanning electron microscope (SEM) and aback-scattered electron image (BEI) was captured at a magnification of×200. Through analysis of four photographs captured from the mold sideto the free surface side with equal intervals by means of an imagegraphic analyzer, the percent volume of α-Fe-containing region was foundto be 1% or less.

Comparative Example 11

The procedure of Example 11 including preparing a raw material andmelt-casting was repeated, except that a roller for casting having asurface roughness, as represented by 10-point average roughness (Rz), of3.0 μm was employed.

The thus-produced alloy flakes were evaluated in a manner similar tothat of Example 11. The alloy flakes were found to have a surface (moldside) roughness, as represented by 10-point average roughness (Rz), of3.4 μm and to have a percent volume of α-Fe-containing region of 8%.

Working examples of production of rare earth magnets will next bedescribed.

Example 13

The flakes of the main phase alloy produced in Example 11 were subjectedto hydrogen decrepitation. Hydrogen absorption step—the step precedinghydrogen decrepitation—was performed under the conditions: 100% hydrogenatmosphere, 2 atm, and retention time of 1 hour. The temperature of thealloy flakes at the start of hydrogen absorption reaction was 25° C.Hydrogen desorption step—subsequent step—was performed under theconditions: vacuum of 0.133 hPa, 500° C., and retention time of 1 hour.To the powder produced through hydrogen decrepitation, zinc stearatepowder was added in an amount of 0.07% by mass. The mixture wassufficiently mixed in a 100% nitrogen atmosphere by use of a V-blender,and then micro-pulverized by use of a jet mill in a nitrogen atmosphereincorporated with oxygen (4,000 ppm). The resultant powder wassufficiently mixed again in a 100% nitrogen atmosphere by use of aV-blender. The obtained powder was found to have an oxygen concentrationof 1,800 ppm. Through analysis of the carbon concentration of thepowder, the zinc stearate content of the powder was calculated to be0.05% by mass.

The boundary phase alloy was prepared in the following manner.Neodymium, dysprosium, ferroboron, aluminum, cobalt, copper, and ironwere mixed to thereby obtain the following alloy composition: Nd: 35.0%by mass; Dy: 20% by mass; B: 0.70% by mass; Al: 0.30% by mass; Co: 25.0%by mass; Cu: 1.00% by mass, and a balance of iron. The resulting mixturewas melted in an alumina crucible in an argon gas atmosphere (1 atm) byuse of a high-frequency induction melting furnace. The resulting moltenalloy was cast through centrifugal casting. During casting, the moltenalloy was deposited on the inner wall of the mold at an averagedeposition rate of 0.03 cm/s. The rotation rate of the mold was modifiedsuch that centrifugal force is adjusted to 20G The thus-produced alloyflakes were found to have a thickness of 8 to 11 mm.

In a manner similar to the case of the main phase alloy, the boundaryphase alloy was subjected to hydrogen decrepitation,micro-pulverization, and mixing. The obtained powder was found to havean oxygen concentration of 3,000 ppm. Through analysis of the carbonconcentration of the powder, the zinc stearate content of the powder wascalculated to be 0.05% by mass.

The aforementioned main phase alloy and the boundary phase alloy weremixed at a ratio by weight of 9:1, and the mixture was sufficientlymixed by use of a V-blender. Subsequently, the thus-obtained powder waspress-molded in a 100% nitrogen atmosphere and a lateral magnetic fieldby use of a molding apparatus. The molding pressure was 1.2 t/cm², andthe magnetic field in the mold cavity was controlled to 15 kOe. Thethus-obtained compact was maintained sequentially in vacuum of 1.33×10⁻⁵hPa at 500° C. for one hour, in vacuum of 1.33×10⁻⁵ hPa at 800° C. fortwo hours, and in vacuum of 1.33×10⁻⁵ hPa at 1,080° C. for two hours forsintering. The density of the sintered product was as sufficiently highas 7.5 g/cm³ or more. The sintered product was further heat-treated at530° C. for one hour in an argon atmosphere.

Magnet characteristics of the thus-produced sintered magnet weremeasured by means of a direct-current BH curve tracer. The results areshown in Table 1. The oxygen content of the main phase alloymicro-powder and the produced sintered magnet are also shown in Table 1.

Example 14

In a manner similar to that employed in Example 13, the main phase alloyflakes obtained in Example 12 were pulverized, to thereby yield apowder. The powder and a micro-powder of the boundary phase alloysimilar to that produced in Example 13 were mixed in a manner similar tothat of Example 13, to thereby produce a rare earth magnet. Magnetcharacteristics of the rare earth magnet produced in Example 14 and theoxygen content of the main phase alloy micro-powder and the producedsintered magnet are also shown in Table 1.

Comparative Example 12

In a manner similar to that employed in Example 13, the main phase alloyflakes obtained in Comparative Example 11 were pulverized, to therebyyield a micro-powder. During pulverization, the rate of pulverization byuse of a jet mill was decreased by 10% (average), as compared with themain phase alloy flakes produced in Example 11. The powder and themicro-powder of the boundary phase alloy produced in Example 13 weremixed in a manner similar to that of Example 13, to thereby produce arare earth magnet. Magnet characteristics of the rare earth magnetproduced in Comparative Example 12 and the oxygen content of the mainphase alloy micro-powder and the produced sintered magnet are also shownin Table 1.

TABLE 1 Oxygen content Magnet of main phase Oxygen micro-powder contentBr iHc (BH)_(max) (ppm) (ppm) (kG) (kOe) (MGOe) Example. 13 1,800 2,20013.3 18.8 41.7 Example. 14 1,900 2,200 13.4 18.3 41.9 Comp. Ex. 12 2,2002,500 13.0 18.7 39.0 Comp. Ex. 13 3,500 3,900 13.2 18.1 40.3

As is clear from Table 1, the rare earth magnet of Comparative Example12 exhibits a smaller residual magnetization as compared with those ofExamples 13 and 14. The small residual magnetization may be attributableto a slight increase in TRE of the micro-powder due to α-Fe which is notpulverized during jet-milling and remains in the jet mill.

Comparative Example 13

A rare earth magnet having a composition similar to that of the magnetobtained in Example 13 was produced through the single-alloy method.

The procedure of Example 11 including strip casting was repeated, exceptthat the alloy composition was replaced by the following composition:TRE: 31.15% by mass (in R, Nd: 52.4% by mass, Pr: 41.2% by mass, and Dy:6.4% by mass); B: 0.97% by mass; Al: 0.30% by mass; Co: 2.50% by mass,Cu: 0.10% by mass, and a balance of iron, to thereby produce alloyflakes.

In a manner similar to that of Example 13, the alloy flakes werepulverized, to thereby yield a micro-powder, and a magnet was producedonly by use of the micro-powder as a magnet source. The oxygen contentand magnet characteristics of the produced magnet are shown in Table 1.Analysis of the magnet produced in Comparative Example 13 has revealedthat the difference in composition between this magnet and the magnetproduced in Example 13 falls within a range of an analytical error.

As shown in Table 1, the rare earth magnet of Comparative Example 13 hasa higher oxygen content and exhibits a smaller residual magnetization ascompared with those of Examples 13 and 14. The above properties may beattributable to oxidation of micro-powder during magnet production stepsand difference in orientation feature provided during orientation in amagnetic field.

Consequently, the first aspect of the present invention provides analloy for a rare earth magnet in which formation of α-Fe is prevented,even when Nd is partially substituted by Pr which is advantageous bothfor cost and characteristics. When employed as a low-TRE main phasealloy to be processed by the two-alloy blending method, the alloy servesas a remarkably effective raw material for producing a rare earth magnetof excellent magnet characteristics.

In addition, when a rare earth magnet is produced by use of thePr-containing main phase alloy for a rare earth magnet of the presentinvention and the boundary phase alloy having a low Pr content, Pr isprovided from the main phase alloy to the magnet. Thus, the invention,which overcomes drawbacks of high-Pr-content rare earth magnets, canprovide a magnet which has improved resistance to oxidation causedduring production step, improved corrosion resistance, and improvedorientation feature in a magnetic field.

Second Aspect

FIG. 7 shows a back-scattered electron image, observed under an SEM(scanning electron microscope), of a cross-section of an Nd—Fe—B alloy(Nd: 31.5 mass %) flake which has been cast through a conventional SCmethod. In FIG. 7, the left side corresponds to the mold side, and theright side to the free surface side. On the mold side, the alloy flakehas a surface roughness, as represented by 10-point average roughness(Rz), of 3.4 μm.

In FIG. 7, white areas correspond to Nd-rich phase (R-rich phase iscalled Nd-rich phase, since R consists of Nd, here). From the centerportion to the free surface side (the surface opposite the mold side) ofthe alloy flake, the Nd-rich phase assumes the form of lamellar portionsextending in the thickness direction, or the form of a small pool oforiented lamellar fragments. In contrast, the Nd-rich phase on the moldside assumes very minute granular form as compared with other portions,and such granular Nd-rich phases are dispersed at random in a region onthe mold side. The present inventor denominates such a region “fineR-rich phase region” (when R predominantly comprises Nd, the region iscalled fine Nd-rich phase region) and distinguishes this region fromother regions. The fine R-rich phase region is generally formed from themold side and extends to the center portion. A portion from the centerto the free surface side where no fine R-rich phase region is present iscalled a “normal portion.”

During hydrogen decrepitation of R-T-B alloy flakes for producing asintered magnet, the volume of R-rich phase increases by absorbinghydrogen, thereby forming a fragile hydride. Thus, when hydrogendecrepitation is performed, microcracks are formed along or from theR-rich phase contained in the alloy. In the subsequentmicro-pulverization step, the alloy flakes are crushed by virtue of alarge amount of microcracks generated in hydrogen decrepitation.Therefore, when the R-rich phase is dispersed more finely in the alloy,the particle size of the resultant micro-powder tends to be smaller.Thus, as compared with a normal portion, the fine R-rich phase region isreadily crushed to form minute particles. For example, the alloy powderobtained from a normal portion has an average particle size of about 3μm as measured by means of FSSS (Fisher Sub-Sieve Sizer), whereas thealloy powder obtained from the fine R-rich phase region contains a largeportion of micropowder having a particle size of 1 μm or less, resultingin a broad particle size distribution profile of the micropluverizedproduct.

Japanese Patent Application Laid-Open (kokai) Nos. 09-170055 and10-36949 disclose that the dispersion state of R-rich phase in an R-T-Balloy can be controlled by regulating a cooling rate of molten alloysolidified during casting or by heat treatment. However, in contrast tothe case of a normal portion, behavior of the R-rich phase present inthe fine R-rich phase region is difficult to control by regulating acooling rate of solidified molten metal or by heat treatment, and theR-rich phase is not widely dispersed but remains finely dispersed.

The percent volume of the fine R-rich phase region can be determined inthe following manner. FIG. 9 is a back-scattered electron image of thesame observation area as that of FIG. 7, but in FIG. 9 the boundarybetween the fine R-rich phase region and the normal portion is specifiedby the line. Since the boundary between two regions can be readilyidentified through observation of the dispersion state of R-rich phase,the percent area of the fine R-rich phase region in the observation areacan be calculated by means of a graphic image analyzer. The percent areain the cross-section corresponds to the percent volume of the alloy.Upon measurement of percent volume of fine R-rich phase region, the fineR-rich phase region content greatly varies among alloy flakes or withinone alloy flake, even when the alloy flakes are cast simultaneously.Thus, graphic image analysis is performed by use of about 5 to about 10flakes under a wide observation area at a low magnification of about 50to about 100 times, and obtained percent area values are averaged, tothereby calculate the percent volume of the fine R-rich phase region forthe entirety of the alloy.

FIG. 8 is a back-scattered electron image of a cross-section of an R-T-Balloy flake (Nd: 31.5 mass %) falling within a scope of the presentinvention. In FIG. 8, the left side corresponds to the mold side and theright side to the free surface side. The alloy flake of the presentinvention is characterized in that formation of fine R-rich phase regionis suppressed by means of controlling the roughness of the mold sidesurface of the flake produced through strip casting. As shown in FIG. 8,the alloy flake of the present invention contains no fine R-rich phaseregion on the mold side, and R-rich phase is dispersed, from the moldside to the free surface side, with remarkably excellent uniformity.

The relationship between the fine R-rich phase region and the surfaceroughness of the mold side surface of an alloy flake produced throughthe strip casting method can be described as follows.

In order to obtain a smooth mold side surface of an alloy flake, thesurface of a rotating roller for casting must be smooth and have highwettability with respect to the molten alloy. When such a rotatingroller is employed, heat is transferred from the molten alloy to themold at remarkably high efficiency (i.e., heat transfer coefficient ishigh). Thus, the mold side alloy is rapidly cooled excessively. The fineR-rich phase region is considered to be highly prone to be generatedthrough excessively rapid cooling of the portion of the alloy on themold side resulting from the large heat transfer coefficient of themolten alloy to the mold.

In contrast, when the surface of the rotating roller for casting isfinely roughened, the minute irregularities formed on the surface of therotating roller for casting cannot be filled completely with the moltenalloy, because of its viscosity. Thus, a portion of the alloy remainsnot in contact with the roller, thereby lowering the heat transfercoefficient. As a result, a portion of the alloy on the mold side is notrapidly cooled to an excessive extent. Accordingly, the above mechanismis considered to prevent generation of the fine R-rich phase region.

When the surface roughness of the rotating roller for casting increases,the surface roughness of the mold side surface of the alloy flakenecessarily increases, through transfer of the irregularities of theroller to the mold side surface of the alloy flake to some extent. Thus,prevention of excessive heat transfer during solidification of themolten alloy, as described above, is considered to be the reason whygeneration of R-rich phase in an alloy flake having an appropriatesurface roughness on the mold side is prevented.

However, when the surface roughness of the rotating roller for castingincreases excessively, the irregularities can be filled with the moltenalloy, thereby increasing heat transfer coefficient and furtherincreasing the surface roughness of the mold side of the produced alloyflake. In this case, percent volume of the fine R-rich phase regionincreases.

Even when a conventional SC method is employed, the produced alloyflakes include, to some extent, those having a uniform microstructure asshown in FIG. 8. However, alloy flakes having large portions of fineR-rich phase regions as shown in FIG. 7 are also producedsimultaneously, thereby deteriorating uniformity in the entiremicrostructure of the resultant alloy. Failure to attain uniformity inmicrostructure of the alloy produced through a conventional SC method isattributable to difference in conditions of contact between the rollersurface and the molten alloy; e.g., the fine surface state of therotating roller for casting, molten alloy supply conditions, and theatmosphere during casting.

In contrast, the rotating roller for casting according to the presentinvention is imparted with appropriate surface roughness. Thus,excessive heat transfer during solidification of molten alloy isprevented, to thereby suppress, at high reproducibility, generation offine R-rich phase region. As a result, alloy flakes having such auniform microstructure as shown in FIG. 8 can be produced at high yield.

The second aspect of the present invention will next be described indetail.

(21) Strip Casting Method

The present invention is drawn to an R-T-B alloy flake for rare earthmagnets which is produced through the strip casting method. Herein,casting of R-T-B alloy through the strip casting method will bedescribed.

FIG. 5 is a schematic view showing a casting apparatus employed in stripcasting. Generally, when an R-T-B alloy is cast, the alloy is mademolten by use of a refractory crucible 1 in vacuum or an inert gasatmosphere, because it is highly active. The thus-molten alloy ismaintained at 1,350 to 1,500° C. for a predetermined period of time, andsupplied, via a tundish 2 having optional flow-control means orslag-removing means, to a rotating roller 3 for casting whose interioris cooled with water. The rate of supplying the molten alloy and therotation speed of the rotating roller are appropriately regulated inaccordance with the thickness of the alloy flakes to be produced.Generally, the rotation speed of the rotating roller is about 1 to about3 m/s (in terms of peripheral velocity). The rotating roller for castingis preferably made of copper or copper alloy, from the viewpoint of highthermal conductivity and availability. The surface of the rotatingroller for casting is prone to adsorb metallic material, depending onthe material and surface conditions of the rotating roller. Thus,provision of an optional cleaning apparatus stabilizes qualities of thecast R-T-B alloy. The alloy 4 solidified on the rotating roller isreleased from the roller on the side opposite the tundish side andcollected into a collection container 5. The microstructure of R-richphase present in the normal portion can be controlled by means ofheating/cooling means provided in the collection container.

The alloy flake of the present invention preferably has a thickness ofat least 0.1 mm and not greater than 0.5 mm. When the thickness of thealloy flake is less than 0.1 mm, solidification rate increasesexcessively, thereby providing an excessively small crystal grain size,which can be equivalent to the particle size of micro-pulverized powderapplied to the magnet production step. In this case, percent orientationand magnetization of the produced magnets are problematicallydeteriorated. A thickness of the alloy flake in excess of 0.5 mm resultsin problems, such as deterioration of Nd-rich phase dispersibilitystemming from a decrease in solidification rate, and problematicprecipitation of α-Fe.

(22) Surface Roughness of the Cast Surface of the Rotating Roller forCasting

According to the present invention, when an R-T-B magnet alloy is castthrough a strip casting method, the surface roughness, as represented by10-point average roughness (Rz), of the cast surface of a rotatingroller for casting is controlled to fall within a range of 5 μm to 100μm.

Herein, the term “surface roughness” refers to a surface roughnessdetermined under the conditions specified in JIS B 0601 “Surfaceroughness—Definitions and Designation,” and 10-point average roughness(Rz) is defined therein. Specifically, a surface to be measured is cutwith a plane which is perpendicular thereto, to thereby obtain a contourappearing on a cut end (profile curve). Any surface waviness componentlonger than a prescribed wavelength is cut off from the profile curve bymeans of a phase-compensation-type high-pass filter or a similar device,to thereby obtain a curve (roughness curve). Only the reference lengthis sampled from the roughness curve in the direction of its mean line,and the sum of the average value of absolute values of the heights ofthe five highest profile peaks (Yp) and the depths of the five deepestprofile valleys (Yv) measured in the vertical direction from the meanline of this sampled portion is calculated, to thereby obtain the10-point average roughness (Rz). Measurement parameters such asreference length are defined in the above JIS B 0601, as standard valuesof reference length for determining corresponding surface roughnessvalues.

Since the surface roughness of the mold side of an alloy flake oftenvaries in a wide range among samples to be measured, an average value ofsurface roughness for at least five flakes should be employed.

When the surface roughness is 5 μm or less, the effect of irregularitiesprovided on the surface of the rotating roller for casting cannot beattained, thereby providing a large area of contact between the moltenalloy and the roller and increasing the heat transfer coefficient. Thus,fine R-rich phase region is easily formed. When the surface roughness is5 μm or more, the minute irregularities formed on the surface of therotating roller cannot be completely filled with the molten alloy,because of its viscosity. Thus, many portions of the alloy remain not incontact with the roller, thereby lowering the heat transfer coefficient.As a result, formation of fine R-rich phase in the alloy can beprevented. The surface roughness, as represented by 10-point averageroughness (Rz), is preferably at least 10 μm.

When the surface roughness of the rotating roller for casting is inexcess of 100 μm, interspacing between peaks (or valleys) generallyincreases with the increase of the depth of the irregularities of therotating roller. Thus, the molten alloy can enter cavities formed on therotating roller, and the heat transfer coefficient readily increasesexcessively, thereby readily forming fine R-rich phase region in thealloy. Therefore, the surface roughness of the rotating roller forcasting is regulated to be 100 μm or less, preferably 50 μm or less.

(23) Surface Roughness of R-T-B Alloy Flakes

According to the present invention, at least one surface of the R-T-Balloy flake for rare earth magnets has a surface roughness, asrepresented by 10-point average roughness (Rz), falling within a rangeof 5 μm to 50 μm. The side on which the above roughness is provided isthe mold side where solidification starts during strip casting, and thesurface roughness of the rotating roller is transferred to the moldside. As mentioned above, when the surface roughness of the mold side is5 μm or less or at least 50 μm, percent volume of the formed fine R-richphase region increases, thereby failing to attain uniformity indispersion state of the R-rich phase in the alloy. As a result, theparticle size distribution profile of the alloy powder micro-pulverizedfor producing sintered magnets becomes broad, thereby deterioratingmagnet characteristics, which is undesirable. Thus, one surface of thealloy flake of the present invention preferably has a surface roughnessfalling within a range of 5 μm to 50 μm, more preferably within a rangeof 7 μm to 25 μm.

(25) Percent Volume of Fine R-Rich Phase Region in the Alloy

According to the present invention, the percent volume of fine R-richphase region in an R-T-B alloy is regulated to 20% or less. Thus, thealloy powder which has been micro-pulverized for producing sinteredmagnets has a sharp particle size distribution profile, thereby yieldingsintered magnets without variation in characteristics.

Method for Producing Rare Earth Sintered Magnet Alloy Powder and Methodfor Producing Rare Earth Sintered Magnets

The rare earth magnet alloy flakes formed of R-T-B alloy which have beencast through the method according to the present invention arepulverized, shaped, and sintered, to thereby produce anisotropicsintered magnets of excellent characteristics.

Typically, pulverization of the alloy flakes is sequentially performedin the order of hydrogen decrepitation and micro-pulverization, tothereby produce an alloy powder having a size of approximately 3 μm(FSSS).

In the present invention, hydrogen decrepitation includes a hydrogenabsorption step as a first step and a hydrogen desorption step as asecond step. In the hydrogen absorption step, hydrogen is caused to beabsorbed predominantly in the R-rich phase of alloy flakes in a hydrogengas atmosphere at 266 hPa to 0.3 MPa·G The R-rich phase is expanded involume due to R hydride generated in this step, to thereby finely breakthe alloy flakes themselves or generate numerous micro-cracks. Hydrogenabsorption is carried out within a temperature range of ambienttemperature to approximately 600° C. However, in order to increaseexpansion in volume of R-rich phase so as to effectively reduce theflakes in size, hydrogen absorption is preferably performed underincreased hydrogen gas pressure and within a temperature range ofambient temperature to approximately 100° C. The time for hydrogenabsorption is preferably one hour or longer. The R hydride formedthrough the hydrogen absorption step is unstable in the atmosphere andreadily oxidized. Thus, the hydrogen-absorbed product is preferablysubjected to hydrogen desorption treatment by maintaining the alloyflakes at about 200 to about 600° C. in vacuum of 1.33 hPa or less.Through this treatment, R hydride can be transformed into a productstable in the atmosphere. The time for hydrogen desorption treatment ispreferably 30 minutes or longer. If the atmosphere is controlled forpreventing oxidation during steps to be carried out after hydrogenabsorption to sintering, hydrogen desorption treatment can also beomitted.

The R-T-B alloy flake produced through the strip casting methodaccording to the present invention is characterized in that R-rich phaseis uniformly dispersed in the alloy flake. The average inter R-richphase spacing, which depends on the particle size of the pulverizedpowder for producing magnets, is preferably 3 μm to 8 μm. Duringhydrogen decrepitation, cracks are introduced to the alloy flake alongor from the R-rich phase therein. Therefore, micro-pulverization of aproduct which has undergone hydrogen decrepitation attains, to a maximumdegree, the effect of the R-rich phase uniformly and finely dispersed inthe alloy, thereby effectively producing an alloy powder exhibiting aremarkably sharp particle size distribution profile. When sinteredmagnets are produced without performing the hydrogen decrepitation step,the produced sintered magnets have poor characteristics (M. Sagawa etal., Proceeding of the 5th international conference on Advancedmaterials, Beijing, China (1999)).

Micro-pulverization is a step of pulverizing R-T-B alloy flakes forattaining a particle size of approximately 3 μm (FSSS). Amongpulverizers for performing the micro-pulverization, a jet mill is mostpreferred, in view of high productivity and a sharp particle sizedistribution profile. By use of alloy flakes according to the presentinvention having a low fine R-rich phase region content, an alloy powderexhibiting a sharp particle size distribution profile can be produced athigh efficiency without variation.

Upon micro-pulverization, the atmosphere is controlled to an inert gasatmosphere such as an argon gas atmosphere or nitrogen gas atmosphere.The inert gas may contain oxygen in an amount of 2% by mass or less,preferably 1% by mass or less. The presence of oxygen enhancespulverization efficiency and attains oxygen concentration of the powderproduced through pulverization to 1,000 to 10,000 ppm, to therebyappropriately stabilize the alloy powder. In addition, abnormal graingrowth during sintering to form magnets can be prevented.

When the alloy powder is molded in a magnetic field, in order to reducefriction between the powder and the inner wall of a mold and to reducefriction generated among powder particles for enhancing orientation, alubricant such as zinc stearate is preferably added to the powder. Theamount of the lubricant to be added is 0.01 to 1% by mass. Although thelubricant may be added before or after micro-pulverization, thelubricant is preferably mixed sufficiently, before molding in magneticfield, in an inert gas atmosphere such as argon gas or nitrogen gas byuse of a mixing apparatus such as a V-blender.

The powder having a particle size of about 3 μm (FSSS) obtained throughmicro-pulverization is press-molded in magnetic field by use of amolding apparatus. The mold to be employed is fabricated from a magneticmaterial and a non-magnetic material in combination in consideration ofthe orientation of magnetic field in the mold cavity. The pressure atmolding is preferably 0.5 to 2 t/cm², and the magnetic field in the moldcavity during molding is preferably 5 to 20 kOe. The atmosphere duringmolding is preferably an inert gas atmosphere such as argon gas ornitrogen gas. However, if the powder has been subjected to theaforementioned anti-oxidation treatment, molding can be performed inair.

Molding may be performed through cold isostatic pressing (CIP) or rubberisostatic pressing (RIP) employing a rubber mold. Since the alloy powderis pressed isostatically through CIP or RIP, variation in orientationduring press-molding is lowered. Thus, the degree of orientation of theproduced compact can be increased as compared with that produced by useof a metal mold, and maximum magnetic energy product can be enhanced.

Sintering of the compact is performed at 1,000 to 1,100° C. Theatmosphere during sintering is preferably an argon gas atmosphere or avacuum atmosphere of 1.33×10⁻² hPa or less. A retention time at thesintering temperature of one hour or longer is preferred. Duringsintering, prior to reaching the sintering temperature, lubricantcontained in the compact and hydrogen contained in the alloy powder mustbe removed as completely as possible from a compact to be sintered. Thelubricant is removed by maintaining the compact preferably under theconditions: in vacuum of 1.33×10⁻² hPa or less or under an argon gasflow atmosphere at reduced pressure; at 300 to 500° C.; and for 30minutes or longer. Hydrogen is removed by maintaining the compactpreferably under the conditions: in vacuum of 1.33×10⁻² hPa or less; at700 to 900° C.; and for 30 minutes or longer.

After completion of sintering, in order to enhance the coercivity ofsintered magnet to be produced, the sintered product may be treated at500 to 650° C. in accordance with needs. An argon gas atmosphere or avacuum atmosphere is preferred, and a retention time of 30 minutes orlonger is preferred.

The rare earth magnet R-T-B alloy flake produced through the methodaccording to the present invention in which formation of fine R-richregion is suppressed can be used suitably for producing bonded magnetsas well as sintered magnets. Production of a bonded magnet by use of therare earth magnet alloy flakes according to the present invention willnext be described.

Firstly, the R-T-B alloy flakes of the present invention undergo heattreatment in advance in accordance with needs. The heat treatment isperformed in order to remove α-Fe contained in the alloy and to coarsencrystal grains. The production steps of the alloy powder for producingbonded magnets includehydrogenation-disproportionation-desorption-recombination (HDDR)treatment. However, α-Fe present in the alloy cannot be removed in theHDDR treatment step, and remaining α-Fe deteriorates magnetism.Therefore, α-Fe must be removed prior to performing the HDDR treatment.

The alloy powder for producing bonded magnets has a mean particle sizeof 50 to 300 μm, which is considerably greater than that of the alloypowder for producing sintered magnets. When the bonded magnet alloyflakes undergo HDDR treatment, crystal orientation of recombined crystalgrains of sub-micron size coincides with crystal orientation of crystalgrains of the starting alloy flakes with a certain range of variance.Thus, when two or more crystal grains having different crystalorientations are contained in each of starting alloy flakes, eachparticle of the bonded magnet alloy powder produced from such alloyflakes will contain crystal grains having different crystalorientations. Thus, the alloy powder includes regions having greatvariance in crystal orientation. In such region, the degree oforientation deteriorates, and maximum magnetic energy product of themagnet is low. In order to avoid such deterioration, the crystal grainscontained in the alloy flakes preferably have a large grain size. Thealloy cast through a rapid-cooling/solidification method (e.g., stripcasting) is prone to have a comparatively small crystal grain size.Thus, coarsening of crystal grains through heat treatment is effectivefor enhancing magnet characteristics.

There are many reports in connection with the method for producing abonded magnet alloy powder through the HDDR method (e.g., T. Takeshitaet al., Proc. 10th Int. Workshop on RE magnets and their application,Kyoto, Vol. 1, P. 551 (1989)). Production of the alloy powder throughthe HDDR method is performed in the following manner.

When R-T-B alloy flakes serving as raw material are heated in a hydrogenatmosphere, the R₂T₁₄B phase—a magnetic phase—decomposes at about 700°C. to about 850° C., to thereby form three phases; i.e., α-Fe, RH₂, andFe₂B. Subsequently, in order to remove hydrogen, the hydrogen atmosphereis replaced by an inert gas atmosphere or a vacuum atmosphere, and thetemperature is maintained approximately in the above range. As a result,separated phases are recombined, to thereby form the R₂T₁₄B phase havingan approximately sub-micron crystal grain size. Upon the above process,if the composition of the alloy or treatment conditions areappropriately modified, the magnetization-easy axis of each recombinedR₂T₁₄B phase (C-axis of R₂T₁₄B phase) is aligned approximately inparallel to the C-axis of R₂T₁₄B phase present in the raw material alloybefore decomposition. Thus, there can be produced an anisotropic magnetpowder in which the magnetization-easy axis of minute crystal grains isaligned.

The alloy which has undergone HDDR treatment is pulverized to form analloy powder having a particle size of about 50 to about 300 μm. By useof the alloy powder, a bonded magnet is produced through a processincluding mixing with resin and press-molding or injection-molding.

Similar to the case of the aforementioned hydrogen decrepitation, fineR-rich phase region is prone to form a micro-powder through HDDRtreatment. Characteristics of the magnetic powder obtained through aHDDR method are deteriorated, as the particle size thereof decreases.Thus, the R-T-B alloy of the present invention in which formation offine R-rich phase is suppressed is suitably used in production a bondedmagnet powder including HDDR treatment.

EXAMPLES Example 21

Neodymium, ferroboron, cobalt, aluminum, copper, and iron were mixed tothereby obtain the following alloy composition: Nd: 31.5% by mass; B:1.00% by mass; Co: 1.0% by mass; Al: 0.30% by mass; Cu: 0.10% by mass;and a balance of iron. The resulting mixture was melted in an aluminacrucible in an argon gas atmosphere (1 atm) by use of a high-frequencyinduction melting furnace. The resulting molten alloy was cast throughstrip casting, to thereby prepare alloy flakes.

The rotating roller for casting having a diameter of 300 mm and made ofpure copper was employed. During casting, the inside of the copperroller was cooled by water. The roller had a cast surface roughness, asrepresented by 10-point average roughness (Rz), of 20 μm and was rotatedat a peripheral velocity of 0.9 m/s, to thereby produce alloy flakeshaving a mean thickness of 0.30 mm.

The thus-produced alloy flakes were found to have a surface (mold side)roughness, as represented by 10-point average roughness (Rz), of 10 μm.Ten flakes were selected from the alloy flakes and polished in a fixedstate. Each flake was observed under a scanning electron microscope(SEM) and a back-scattered electron image (BEI) was captured at amagnification of ×100. Through analysis of the thus-captured photographby means of an image graphic analyzer, the percent volume of fine R-richphase region was found to be 3% or less.

Example 22

The procedure of Example 21 including casting through an SC method wasrepeated, except that a raw material having the following alloycomposition: Nd: 28.5%; B: 1.00% by mass; Co: 1.0% by mass; Al: 0.30% bymass; Cu: 0.10% by mass; and a balance of iron was used, to therebyproduce alloy flakes.

The thus-produced alloy flakes were evaluated in a manner similar tothat of Example 21. The alloy flakes were found to have a surface (moldside) roughness, as represented by 10-point average roughness (Rz), of 9μm and to have a percent volume of fine R-rich phase region of 3% orless.

Comparative Example 21

The procedure of Example 21 including preparing a raw material, melting,and casting through an SC method was repeated, except that a rotatingroller for casting having a surface roughness, as represented by10-point average roughness (Rz), of 3.0 μm was employed, to therebyproduce alloy flakes.

The thus-produced alloy flakes were evaluated in a manner similar tothat of Example 21. The alloy flakes were found to have a surface (moldside) roughness, as represented by 10-point average roughness (Rz), of3.3 μm and to have a percent volume of fine R-rich phase region of 41%.

Comparative Example 22

The procedure of Example 21 including preparing a raw material, melting,and casting through an SC method was repeated, except that a rotatingroller for casting having a surface roughness, as represented by10-point average roughness (Rz), of 120 μm was employed, to therebyproduce alloy flakes.

The thus-produced alloy flakes were evaluated in a manner similar tothat of Example 21. The alloy flakes were found to have a surface (moldside) roughness, as represented by 10-point average roughness (Rz), of86 μm and to have a percent volume of fine R-rich phase region of 29%.

Working examples of production of sintered magnets will next bedescribed.

Example 23

The alloy flakes produced in Example 21 were subjected to hydrogendecrepitation and micro-pulverization by use of a jet mill. Hydrogenabsorption step—the step preceding hydrogen decrepitation—was performedunder the conditions: 100% hydrogen atmosphere, 2 atm, and retentiontime of 1 hour. The temperature of the alloy flakes at the start ofhydrogen absorption reaction was 25° C. Hydrogen desorptionstep—subsequent step—was performed under the conditions: vacuum of 0.133hPa, 500° C., and retention time of 1 hour. To the resultant powder,zinc stearate powder was added in an amount of 0.07% by mass. Themixture was sufficiently mixed in a 100% nitrogen atmosphere by use of aV-blender, and then micro-pulverized by use of a jet mill in a nitrogenatmosphere incorporated with oxygen (4,000 ppm). The resultant powderwas sufficiently mixed again in a 100% nitrogen atmosphere by use of aV-blender. The obtained powder was found to have an oxygen concentrationof 2,500 ppm. Through analysis of the carbon concentration of thepowder, the zinc stearate content of the powder was calculated to be0.05% by mass. The mean particle sizes of the powder, as measured bymeans of a laser diffraction particle size distribution measurementapparatus, were found to be 5.10 μm (D50), 2.10 μm (D10), and 8.62 μm(D90).

Subsequently, the thus-obtained powder was press-molded in a 100%nitrogen atmosphere and a lateral magnetic field by use of a moldingapparatus. The molding pressure was 1.2 t/cm², and the magnetic field inthe mold cavity was controlled to 15 kOe. The thus-obtained compact wasmaintained sequentially in vacuum of 1.33×10⁻⁵ hPa at 500° C. for onehour, in vacuum of 1.33×10⁻⁵ hPa at 800° C. for two hours, and in vacuumof 1.33×10⁻⁵ hPa at 1,050° C. for two hours for sintering. The densityof the sintered product was as sufficiently high as 7.5 g/cm³ or more.The sintered product was further heat-treated at 560° C. for one hour inan argon atmosphere, to thereby produce a sintered magnet.

Magnet characteristics of the sintered magnet were measured by means ofa direct-current BH curve tracer. The results are shown in Table 2. Theoxygen content and particle size of the raw micro-powder for producingthe sintered magnet are also shown in Table 2.

Comparative Examples 23 and 24

In a manner similar to Example 23, alloy flakes produced in ComparativeExamples 21 or 22 were pulverized, to thereby obtain a micro-powder. Theprocedure of molding and sintering performed in Example 23 was repeated,except that the temperature of sintering the micro-powder obtained fromalloy flakes of Comparative Example 21 or 22 was elevated by 20° C. dueto less sinterability of these micro-powders, to thereby produce asintered magnet. Results of evaluation of a sintered magnet producedfrom the alloy flakes of Comparative Example 21 and that produced fromthe alloy flakes of Comparative Example 22 are shown in Table 2 in thecolumns of Comparative Examples 23 and 24, respectively.

Magnet characteristics of the sintered magnets were measured by means ofa direct-current BH curve tracer. The results are shown in Table 2. Theoxygen content and particle size of each raw micro-powder for producingthe sintered magnet are also shown in Table 2.

TABLE 2 Micro-powder Oxygen Magnet content Particle size (μm) Br iHc(BH)_(max) (ppm) D10 D50 D90 (kG) (kOe) (MGOe) Example 23 2,500 2.1 5.18.6 13.6 14.5 44.7 Comp. Ex. 23 3,300 1.6 4.9 8.8 13.5 13.6 43.6 Comp.Ex. 24 3,100 1.8 5.0 8.8 13.6 13.9 44.2 Example 24 — — — —  9.1 13.518.1 Comp. Ex. 25 — — — —  9.1 12.6 17.5

As is clear from Table 2, micro-powders obtained in Comparative Examples23 and 24 have a smaller D10 as compared with that of the micro-powderobtained in Example 23; i.e., contain large amounts of very minuteparticles having a particle size of less than about 1 μm. Since suchminute powders are readily oxidized, micro-powders obtained inComparative Examples 23 and 24 exhibit a slightly higher oxygen contentas compared with that of the micro-powder of Example 23. Magneticcharacteristics of the magnets obtained in Comparative Examples 23 and24 are inferior to those of the magnet of Example 23. The poorcharacteristics are mainly considered to be attributed to coarsening ofcrystal grains, which is caused by increase in sintering temperature by20° C. performed for enhancing sinterability lowered by increase inoxygen content.

Working examples of production of bonded magnets will next be described.

Example 24

The procedure of Example 21 including casting through an SC method wasrepeated, except that a raw material having the following alloycomposition: Nd: 28.5%; B: 1.00% by mass; Co: 10.0% by mass; Ga: 0.5% bymass; and a balance of iron was used, to thereby produce alloy flakes.

The thus-produced alloy flakes were evaluated in a manner similar tothat of Example 21. The alloy flakes were found to have a surface (moldside) roughness, as represented by 10-point average roughness (Rz), of 9μm and to have a percent volume of fine R-rich phase region of 3% orless. The alloy flakes contain no α-Fe.

The above alloy flakes were subjected to HDDR treatment includingannealing in hydrogen (1 atm) at 820° C. for one hour and subsequentannealing in vacuum at 820° C. for one hour. The resultant alloy powderwas pulverized by means of a Brawn mill so as to have a particle size of150 μm or less and blended with an epoxy resin (2.5% by mass). Theresultant mixture was press-formed in a magnetic field of 1.5 T, tothereby obtain a bonded magnet. Magnetic characteristics of the bondedmagnet are shown in Table 2.

Comparative Example 25

The procedure of Comparative Example 21 including melting and castingthrough an SC method was repeated, except that the alloy composition wasreplaced by the alloy composition employed in Example 24, to therebyproduce alloy flakes. The thus-produced alloy flakes were evaluated in amanner similar to that of Example 21. The alloy flakes were found tohave a surface (mold side) roughness, as represented by 10-point averageroughness (Rz), of 3.1 μm and to have a percent volume of fine R-richphase region of 40%.

Subsequently, a bonded magnet was produced in a manner similar to thatof Example 4. Magnetic characteristics of the bonded magnet are shown inTable 2.

As is clear from Table 2, the bonded magnet produced in Example 24exhibits more excellent magnetic characteristics than those of thebonded magnet produced in Comparative Example 25. The bonded magnetproduced in Comparative Example 25 has a high percent volume of fineR-rich phase region and contains a large number of comparatively smallgrains having a grain size of 50 μm or less produced through HDDRtreatment or pulverization. The poor magnetic characteristics areattributable to such a small grain size.

As a result, the alloy flakes according to the present invention, havinga small percent volume of fine R-rich region, exhibit higher uniformityin R-rich phase dispersion state in the alloy as compared withconventional SC materials. Thus, sintered magnets produced from thealloy flakes and bonded magnets produced by use of the flakes through anHDDR method exhibit more excellent magnetic characteristics than thoseof conventional magnets.

Third Aspect

FIG. 7 shows a back-scattered electron image, observed under an SEM(scanning electron microscope), of a cross-section of an Nd—Fe—B alloy(Nd: 31.5 mass %) flake which has been cast through a conventional SCmethod. In FIG. 7, the left side corresponds to the mold side, and theright side to the free surface side. On the mold side, the alloy flakehas a surface roughness, as represented by 10-point average roughness(Rz), of 3.4 μm. The surface is provided with elongated raised/dentedsegments extending in a direction almost in parallel.

In FIG. 7, white areas correspond to Nd-rich phase (R-rich phase iscalled Nd-rich phase, since R is Nd). From the center portion to thefree surface side (the surface opposite the mold side) of the alloyflake, the Nd-rich phase assumes the form of lamellar portions extendingin the thickness direction, or the form of a small pool of orientedlamellar fragments. In contrast, the Nd-rich phase on the mold side hasconsiderably minute grains as compared with other portions, and suchgrains are present at random in a region on the mold side. The presentinventor denominates such a region “fine R-rich phase region” (when Rpredominantly comprises Nd, the region is called fine Nd-rich phaseregion) and distinguishes this region from other regions. The fineR-rich phase region is generally formed from the mold side and extendsto the center portion. A portion from the center to the free surfaceside where no fine R-rich phase region is present is called a “normalportion.”

During hydrogen decrepitation of R-T-B alloy flakes for producing asintered magnet, the volume of R-rich phase increases by absorbinghydrogen, thereby forming a fragile hydride. Thus, when hydrogendecrepitation is performed, microcracks are formed along or from theR-rich phase contained in the alloy. In the subsequentmicro-pulverization step, the alloy flakes are crushed by virtue of alarge amount of microcracks generated in hydrogen decrepitation.Therefore, when the R-rich phase is dispersed more finely in the alloy,the particle size of the resultant micro-powder tends to be smaller.Thus, as compared with a normal portion, the fine R-rich phase region isreadily crushed to form minute particles. For example, the alloy powderobtained from a normal portion has an average particle size of about 3μm as measured by means of FSSS (Fisher Sub-Sieve Sizer), whereas thealloy powder obtained from the fine R-rich phase region contains a largeportion of micro-powder having a particle size of 1 μm or less,resulting in a broad particle size distribution profile of themicro-pulverized product.

Japanese Patent Application Laid-Open (kokai) Nos. 09-170055 and10-36949 disclose that the dispersion state of R-rich phase in an R-T-Balloy can be controlled by regulating a cooling rate of molten alloysolidified during casting or by heat treatment. However, in contrast tothe case of a normal portion, behavior of the R-rich phase present inthe fine R-rich phase region is difficult to control by regulating acooling rate of solidified molten metal or by heat treatment, and theR-rich phase is not widely dispersed but remains finely dispersed.

The percent volume of the fine R-rich phase region can be determined inthe following manner. FIG. 9 is a back-scattered electron image of thesame observation area as that of FIG. 7, but in FIG. 9 the boundarybetween the fine R-rich phase region and the normal portion is specifiedby the line. Since the boundary between two regions can be readilyidentified through observation of the dispersion state of R-rich phase,the percent area of the fine R-rich phase region in the observation areacan be calculated by means of a graphic image analyzer. The percent areain the cross-section corresponds to the percent volume of the alloy.Upon measurement of percent volume of fine R-rich phase region, the fineR-rich phase region content greatly varies among alloy flakes or withinone alloy flake, even when the alloy flakes are cast simultaneously.Thus, graphic image analysis is performed by use of about 5 to about 10flakes under a wide observation area at a low magnification of about 50to about 100 times, and obtained percent area values are averaged, tothereby calculate the percent volume of the fine R-rich phase region forthe entirety of the alloy.

FIG. 8 is a back-scattered electron image of a cross-section of an R-T-Balloy flake (Nd: 31.5 mass %) falling within a scope of the Third Aspectof the present invention. In FIG. 8, the left side corresponds to themold side and the right side to the free surface side. The alloy flakeof the Third Aspect of the present invention is characterized in thatformation of fine R-rich phase region is suppressed by means ofcontrolling the roughness of the mold side surface of the flake producedthrough strip casting and by forming on the surface elongatedraised/dented segments so as to cross one another. The alloy flake shownin FIG. 8 has a mold side roughness of 3.2 μm, which is approximatelyequal to that of the alloy flake shown in FIG. 7. However, the alloyflake of the present invention contains no fine R-rich phase region onthe mold side, and R-rich phase is dispersed, from the mold side to thefree surface side, with remarkably excellent uniformity.

The relationship between the fine R-rich phase region and the surfaceroughness of the mold side surface of an alloy flake produced throughthe strip casting method can be described as follows.

In order to obtain a smooth mold side surface of an alloy flake, thesurface of a rotating roller for casting must be smooth and have highwettability with respect to the molten alloy. When such a rotatingroller is employed, heat is transferred from the molten alloy to themold at remarkably high efficiency (i.e., heat transfer coefficient ishigh). Thus, the mold side alloy is rapidly cooled excessively. The fineR-rich phase region is considered to be highly prone to be generatedthrough excessively rapid cooling of the portion of the alloy on themold side resulting from the large heat transfer coefficient of themolten alloy to the mold.

In contrast, when the surface of the rotating roller for casting isminutely roughened, the minute irregularities formed on the surface ofthe rotating roller for casting cannot be filled completely with themolten alloy, because of its viscosity. Thus, a portion of the alloyremains not in contact with the roller, thereby lowering the heattransfer coefficient. As a result, a portion of the alloy on the moldside is not rapidly cooled to an excessive extent. Accordingly, theabove mechanism is considered to prevent generation of the fine R-richphase region. When the surface roughness of the rotating roller forcasting increases, the surface roughness of the mold side surface of thealloy flake necessarily increases, through transfer of theirregularities of the roller to the mold side surface of the alloy flaketo some extent. Thus, prevention of excessive heat transfer duringsolidification of the molten alloy, as described above, is considered tobe the reason why generation of R-rich phase in an alloy flake having anappropriate surface roughness on the mold side is prevented.

In connection with the morphology of raised/dented segments, when thesesegments are elongated segments which are not crossing one another, eachof contact and non-contact portions between the molten alloy and theroller tends to extend along an elongated raised/dented segment.Accordingly, the internal microstructure is also prone to exhibitcontinuity along such a raised/dented segment. In this case, if a fineR-rich phase region is formed in an elongated raised/dented segment forsome reason, there arises the risk of growth of the fine R-rich phaseregion in the entire portion of the elongated raised/dented segment.

However, when elongated raised/dented segments cross one another, thesegments on the surface are fragmented, and the continuity of theinternal microstructure of the alloy is cut at crossing points.Furthermore, an elongated raised segment is necessarily cut at acrossing point by a linear dented segment. At the raised segment,contact area between the molten alloy and the surface of the roller forcasting increases, thereby promoting heat transfer. Thus, fine R-richphase region is considered to be readily formed through rapidcooling/solidification. However, such a fragmented segment preventsextension of fine R-rich phase, even if the fine R-rich phase region isformed.

According to the method for producing a rare-earth-containing alloyflake including a strip casting method, a rotating roller for casting isemployed, the roller having, on the cast surface, a plurality ofelongated raised/dented segments formed so as to cross one another andhaving a surface roughness of the cast surface, as represented by10-point average roughness (Rz), falling within a range of 3 μm to 30μm. The method can provide a rare-earth-containing alloy flake, whereinat least one surface of the alloy flake has a plurality of elongatedraised/dented segments formed so as to cross one another; and thesurface having the elongated raised/dented segments has a surfaceroughness, as represented by 10-point average roughness (Rz), fallingwithin a range of 3 μm to 30 μm. According to the Third Aspect of thepresent invention, formation of fine R-rich phase region is prevented,thereby attaining a uniform microstructure, even though the surfaceroughness is small compared with the case of the second Aspect of thepresent invention. In addition, since a small surface roughness of therotating roller for casting decreases the amount of grind for regulatingthe roller surface, the service life of the rotating roller for castingcan be prolonged. According to the Third Aspect of the presentinvention, standards for controlling surface conditions of the rollercan be simplified, since effects exerted by surface roughness becomesmaller.

Even when a conventional SC method is employed, the produced alloyflakes include, to some extent, those having a uniform microstructure asshown in FIG. 8. However, alloy flakes having large portions of fineR-rich phase regions as shown in FIG. 7 are also producedsimultaneously, thereby deteriorating uniformity in the entiremicrostructure of the resultant alloy. Failure to attain uniformity inmicrostructure of the alloy produced through a conventional SC methodmay be attributable to difference in conditions of contact between theroller surface and the molten alloy; e.g., the fine surface state of therotating roller for casting, molten alloy supply conditions, and theatmosphere during casting. Surface irregularity provided on the surfaceof a rotating roller for casting prevents excessive heat transfer duringsolidification of molten alloy, to thereby suppress, at highreproducibility, generation of fine R-rich phase region.

In addition, according to the Third Aspect of the present invention,elongated raised/dented segments which cross one another are provided onthe surface of a rotating roller for casting. Thus, effect of preventingformation of fine R-rich phase region is strengthened and issatisfactory, even when the surface roughness is comparatively small. Asa result, alloy flakes having such a uniform microstructure as shown inFIG. 8 can be produced at high yield.

The present invention will next be described in detail.

(31) Strip Casting (SC) Method

The present invention is drawn to a rare-earth-containing alloy flakewhich is produced through the strip casting method. Herein, casting ofR-T-B alloy through the strip casting method will be described.

FIG. 4 is a schematic view showing a casting apparatus employed in stripcasting. Generally, when an R-T-B alloy is cast, the alloy is mademolten by use of a refractory crucible 1 in vacuum or an inert gasatmosphere, because it is highly active. The thus-molten alloy ismaintained at 1,350 to 1,500° C. for a predetermined period of time, andsupplied, via a tundish 2 having optional flow-control means orslag-removing means, to a rotating roller 3 for casting whose interioris cooled with water. The rate of supplying the molten alloy and therotation speed of the rotating roller are appropriately regulated inaccordance with the thickness of the alloy flakes to be produced.Generally, the rotation speed of the rotating roller is about 1 to about3 m/s (in terms of peripheral velocity). The rotating roller for castingis preferably made of copper or copper alloy, from the viewpoint of highthermal conductivity and availability. The surface of the rotatingroller for casting is prone to adsorb metallic material, depending onthe material and surface conditions of the rotating roller. Thus,provision of an optional cleaning apparatus stabilizes qualities of thecast R-T-B alloy. The alloy 4 solidified on the rotating roller isreleased from the roller on the side opposite the tundish side andcollected into a collection container 5. The microstructure of R-richphase present in the normal portion can be controlled by means ofheating/cooling means provided in the collection container.

The alloy flake of the present invention preferably has a thickness ofat least 0.1 mm and not greater than 0.5 mm. When the thickness of thealloy flake is less than 0.1 mm, solidification rate increasesexcessively, thereby providing an excessively small crystal grain size,which is equivalent to the particle size of micro-pulverized powderapplied to the magnet production step. In this case, percent orientationand magnetization of the produced magnets are problematicallydeteriorated. A thickness of the alloy flake in excess of 0.5 mm resultsin problems, such as deterioration of Nd-rich phase dispersibilitystemming from a decrease in solidification rate, and problematicprecipitation of α-Fe.

(32) Surface Roughness of the Cast Surface of the Rotating Roller forCasting

According to the Third Aspect of the present invention, when an R-T-Bmagnet alloy is cast through a strip casting method, the surfaceroughness, as represented by 10-point average roughness (Rz), of thecast surface of a rotating roller for casting is controlled to fallwithin a range of 3 μm to 30 μm.

Herein, the term “surface roughness” refers to a surface roughnessdetermined under the conditions specified in JIS B 0601 “Surfaceroughness—Definitions and Designation,” and 10-point average roughness(Rz) is defined therein. Specifically, a surface to be measured is cutwith a plane which is perpendicular thereto, to thereby obtain a contourappearing on a cut end (profile curve). Any surface waviness componentlonger than a prescribed wavelength is cut off from the profile curve bymeans of a phase-compensation-type high-pass filter or a similar device,to thereby obtain a curve (roughness curve). Only the reference lengthis sampled from the roughness curve in the direction of its mean line,and the sum of the average value of absolute values of the heights ofthe five highest profile peaks (Yp) and the depths of the five deepestprofile valleys (Yv) measured in the vertical direction from the meanline of this sampled portion is calculated, to thereby obtain the10-point average roughness (Rz). Measurement parameters such asreference length are defined in the above JIS B 0601, as standard valuesof reference length for determining corresponding surface roughnessvalues.

Since the surface roughness of the mold side of an alloy flake oftenvaries in a wide range among samples to be measured, an average value ofsurface roughness for at least five flakes should be employed.

(33) Morphology of Surface Irregularity of the Cast Surface of aRotating Roller for Casting

According to the Third Aspect of the present invention, surfaceirregularities of the cast surface are generally provided by a pluralityof elongated raised/dented segments formed on the cast surface so as tocross one another.

When these segments are in line form, each of contact and non-contactportions between the molten alloy and the roller tends to extend alongan elongated raised/dented segment. Accordingly, the internalmicrostructure is also prone to exhibit continuity along such araised/dented segment. In this case, if a fine R-rich phase region isformed in an elongated raised/dented segment for some reason, therearises the risk of growth of the fine R-rich phase region in the entireportion of the elongated raised/dented segment.

However, when elongated raised/dented segments cross one another, thesegments on the surface are fragmented, and the continuity of theinternal microstructure of the alloy is cut at crossing points. Thus,even though the fine R-rich phase region is formed, extension of fineR-rich phase can be prevented.

According to the Third Aspect of the present invention, uniformmicrostructure can be provided through effect of elongated raised/dentedsegments provided so as to cross one another, even though the surfaceroughness, as represented by 10-point average roughness (Rz), iscomparatively small (i.e., falling within a range of 3 μm to 30 μm).

However, when the surface roughness is 3 μm or less, effect exerted bythe presence of irregularities is unsatisfactory. Thus, heat transfer ispromoted through increased contact between the molten alloy and thesurface of a rotating roller for casting, thereby readily forming fineR-rich phase region in the alloy.

When the surface roughness of the rotating roller for casting is inexcess of 30 μm, a solidified alloy flake is engaged with the rollersurface and difficult to peel from the roller, thereby possibly causingtrouble such as breakage of a tundish. Therefore, the surface roughnessof the rotating roller for casting is controlled to 30 μm or less.

(34) Surface Roughness of Rare-Earth-Containing Alloy Flakes andMorphology of Irregularities

According to the Third Aspect of the present invention, at least onesurface of the rare-earth-containing alloy flake has a surfaceroughness, as represented by 10-point average roughness (Rz), fallingwithin a range of 3 μm to 30 μm. The surface roughness is generallyprovided by a plurality of elongated raised/dented segments formed onthe surface so as to cross one another.

The side on which irregularities of the above roughness are formed isthe mold side where solidification starts during strip casting, and thesurface irregularities of the rotating roller are transferred to themold side. As mentioned above, when the surface roughness of the moldside is 3 μm or less, percent volume of the formed fine R-rich phaseregion increases, thereby failing to attain uniformity in dispersionstate of the R-rich phase in the alloy. As a result, the particle sizedistribution profile of the alloy powder micro-pulverized for producingsintered magnets becomes broad, thereby deteriorating magnetcharacteristics, which is undesirable. When the surface roughness is 30μm or more, trouble occurs readily in the course of casting of thealloy.

Thus, one surface of the alloy flake of the Third Aspect of the presentinvention preferably has a surface roughness falling within a range of 3μm to 30 μm.

(35) Percent Volume of Fine R-Rich Phase Region in the Alloy

According to the present invention, the percent volume of fine R-richphase region in an R-T-B alloy is regulated to 20% or less. Thus, thealloy powder which has been micro-pulverized for producing sinteredmagnets has a sharp particle size distribution profile, thereby yieldingsintered magnets without variation in characteristics.

(36) Method for Producing Rare Earth Sintered Magnet Alloy Powder andMethod for Producing Rare Earth Sintered Magnets

The rare earth magnet alloy flakes formed of R-T-B alloy for producing amagnet which flakes have been cast through the method according to thepresent invention are pulverized, shaped, and sintered, to therebyproduce anisotropic sintered magnets of excellent characteristics.

Typically, pulverization of the alloy flakes is sequentially performedin the order of hydrogen decrepitation and micro-pulverization, tothereby produce an alloy powder having a size of approximately 3 μm(FSSS). In the present invention, hydrogen decrepitation includes ahydrogen absorption step as a first step and a hydrogen desorption stepas a second step. In the hydrogen absorption step, hydrogen is caused tobe absorbed predominantly in the R-rich phase of alloy flakes in ahydrogen gas atmosphere at 266 hPa to 0.3 MPa. The R-rich phase isexpanded in volume due to R hydride generated in this step, to therebyminutely break the alloy flakes themselves or generate numerousmicro-cracks. Hydrogen absorption is carried out within a temperaturerange of ambient temperature to approximately 600° C. However, in orderto increase expansion in volume of R-rich phase so as to effectivelyreduce the flakes in size, hydrogen absorption is preferably performedunder increased hydrogen gas pressure and within a temperature range ofambient temperature to approximately 100° C. The time for hydrogenabsorption is preferably one hour or longer. The R hydride formedthrough the hydrogen absorption step is unstable in the atmosphere andreadily oxidized. Thus, the hydrogen-absorbed product is preferablysubjected to hydrogen desorption treatment by maintaining the alloyflakes at about 200 to about 600° C. in vacuum of 1.33 hPa or less.Through this treatment, R hydride can be transformed into a productstable in the atmosphere. The time for hydrogen desorption treatment ispreferably 30 minutes or longer. If the atmosphere is controlled forpreventing oxidation during steps to be carried out after hydrogenabsorption to sintering, hydrogen desorption treatment can also beomitted.

The R-T-B alloy flake produced through the strip casting methodaccording to the present invention is characterized in that R-rich phaseis uniformly dispersed in the alloy flake. The average inter R-richphase spacing, which depends on the particle size of the pulverizedpowder for producing magnets, is preferably 3 μm to 8 μm. Duringhydrogen decrepitation, cracks are introduced to the alloy flake alongor from the R-rich phase therein. Therefore, micro-pulverization of aproduct which has undergone hydrogen decrepitation attains, to a maximumdegree, the effect of the R-rich phase uniformly and finely dispersed inthe alloy, thereby effectively producing an alloy powder exhibiting aremarkably sharp particle size distribution profile. When sinteredmagnets are produced without performing the hydrogen decrepitation step,the produced sintered magnets have poor characteristics (M. Sagawa etal., Proceeding of the 5th international conference on Advancedmaterials, Beijing, China (1999)).

Micro-pulverization is a step of pulverizing R-T-B alloy flakes forattaining a particle size of approximately 3 μm (FSSS). Amongpulverizers for performing the micro-pulverization, a jet mill is mostpreferred, in view of high productivity and a sharp particle sizedistribution profile. By use of alloy flakes according to the presentinvention having a low fine R-rich phase region content, an alloy powderexhibiting a sharp particle size distribution profile can be produced athigh efficiency without variation.

Upon micro-pulverization, the atmosphere is controlled to an inert gasatmosphere such as an argon gas atmosphere or nitrogen gas atmosphere.The inert gas may contain oxygen in an amount of 2% by mass or less,preferably 1% by mass or less. The presence of oxygen enhancespulverization efficiency and attains oxygen concentration of the alloypowder produced through pulverization to 1,000 to 10,000 ppm, to therebyappropriately stabilize the alloy powder. In addition, abnormal graingrowth during sintering to form magnets can be prevented.

When the alloy powder is molded in a magnetic field, in order to reducefriction between the powder and the inner wall of a mold and to reducefriction generated among powder particles for enhancing orientation, alubricant such as zinc stearate is preferably added to the powder. Theamount of the lubricant to be added is 0.01 to 1% by mass. Although thelubricant may be added before or after micro-pulverization, thelubricant is preferably mixed sufficiently, before molding in magneticfield, in an inert gas atmosphere such as argon gas or nitrogen gas byuse of a mixing apparatus such as a V-blender.

The R-T-B alloy powder having a particle size of about 3 μm (FSSS)obtained through micro-pulverization is press-molded in magnetic fieldby use of a molding apparatus. The mold to be employed is fabricatedfrom a magnetic material and a non-magnetic material in combination inconsideration of the orientation of magnetic field in the mold cavity.The pressure at molding is preferably 0.5 to 2 t/cm², and the magneticfield in the mold cavity during molding is preferably 5 to 20 kOe. Theatmosphere during molding is preferably an inert gas atmosphere such asargon gas or nitrogen gas. However, if the powder has been subjected tothe aforementioned anti-oxidation treatment, molding can be performed inair.

Molding may be performed through cold isostatic pressing (CIP) or rubberisostatic pressing (RIP) employing a rubber mold. Since the alloy powderis pressed isostatically through CIP or RIP, variation in orientation ofmagnetization during press-molding is lowered. Thus, the degree oforientation of the produced compact can be increased as compared withthat produced by use of a metal mold, and maximum magnetic energyproduct can be enhanced.

Sintering of the compact is performed at 1,000 to 1,100° C. Theatmosphere during sintering is preferably an argon gas atmosphere or avacuum atmosphere of 1.33×10⁻² hPa or less. A retention time at thesintering temperature of one hour or longer is preferred. Duringsintering, prior to reaching the sintering temperature, a lubricant andhydrogen must be removed as completely as possible from a compact to besintered. The lubricant is removed by maintaining the compact preferablyunder the conditions: in vacuum of 1.33×10⁻² Pa or less or under an Arflow atmosphere at reduced pressure; at 300 to 500° C.; and for 30minutes or longer. Hydrogen is removed by maintaining the compactpreferably under the conditions: in vacuum of 1.33×10⁻² hPa or less; at700 to 900° C.; and for 30 minutes or longer.

After completion of sintering, in order to enhance the coercivity ofsintered magnet to be produced, the sintered product may be treated at500 to 650° C. in accordance with needs. An argon gas atmosphere or avacuum atmosphere is preferred, and a retention time of 30 minutes orlonger is preferred.

The rare earth magnet R-T-B alloy flake produced through the methodaccording to the present invention in which formation of fine R-richregion is suppressed can be used suitably for producing bonded magnetsas well as sintered magnets. Production of a bonded magnet by use of therare earth magnet alloy flakes according to the present invention willnext be described.

Firstly, the R-T-B alloy flakes of the present invention undergo heattreatment in advance in accordance with needs. The heat treatment isperformed in order to remove α-Fe contained in the alloy and to coarsencrystal grains. The production steps of the alloy powder for producingbonded magnets includehydrogenation-disproportionation-desorption-recombination (HDDR)treatment. However, α-Fe present in the alloy cannot be removed in theHDDR treatment step, and remaining α-Fe deteriorates magnetism.Therefore, α-Fe must be removed prior to performing the HDDR treatment.

The alloy powder for producing bonded magnets has a mean particle sizeof 50 to 300 μm, which is considerably greater than that of the alloypowder for producing sintered magnets. When the bonded magnet alloyflakes undergo HDDR treatment, crystal orientation of recombined crystalgrains of sub-micron size coincides with crystal orientation of crystalgrains of the starting alloy flakes with a certain range of variance.Thus, when two or more crystal grains having different crystalorientations are contained in each of starting alloy flakes, eachparticle of the bonded magnet alloy powder produced from such alloyflakes will contain crystal grains having different crystalorientations. Thus, the alloy powder includes regions having greatvariance in crystal orientation. In such region, the degree oforientation deteriorates, and maximum magnetic energy product of themagnet is low. In order to avoid such deterioration, the crystal grainscontained in the alloy flakes preferably have a large grain size. Thealloy cast through a rapid-cooling/solidification method (e.g., stripcasting) is prone to have a comparatively small crystal grain size.Thus, coarsening of crystal grains through heat treatment is effectivefor enhancing magnet characteristics.

There are many reports in connection with the method for producing abonded magnet alloy powder through the HDDR method (e.g., T. Takeshitaet al., Proc. 10th Int. Workshop on RE magnets and their application,Kyoto, Vol. 1, P. 551 (1989)). Production of the alloy powder throughthe HDDR method is performed in the following manner.

When R-T-B alloy flakes serving as raw material are heated in a hydrogenatmosphere, the R₂T₁₄B phase—a magnetic phase—decomposes at about 700°C. to about 850° C., to thereby form three phases; i.e., α-Fe, RH₂, andFe₂B. Subsequently, in order to remove hydrogen, the hydrogen atmosphereis replaced by an inert gas atmosphere or a vacuum atmosphere, and thetemperature is maintained approximately in the above range. As a result,separated phases are recombined, to thereby form the R₂T₁₄B phase havingan approximately sub-micron crystal grain size. Upon the above process,if the composition of the alloy or treatment conditions areappropriately modified, the magnetization-easy axis of each recombinedR₂T₁₄B phase (C-axis of R₂T₁₄B phase) is aligned approximately inparallel to the C-axis of R₂T₁₄B phase present in the raw material alloybefore decomposition. Thus, there can be produced an anisotropic magnetpowder in which the magnetization-easy axis of minute crystal grains isaligned.

The alloy which has undergone HDDR treatment is pulverized to form analloy powder having a particle size of about 50 to about 300 μm. By useof the alloy powder, a bonded magnet is produced through a processincluding mixing with resin and press-molding or injection-molding.

Similar to the case of the aforementioned hydrogen decrepitation, fineR-rich phase region is prone to form a micro-powder through HDDRtreatment. Characteristics of the magnetic powder obtained through aHDDR method are deteriorated, as the particle size thereof decreases.Thus, the R-T-B alloy of the present invention in which formation offine R-rich phase is suppressed is suitably used in production a bondedmagnet powder including HDDR treatment.

Recently, it has been reported that surface roughness parameters (Sm/Raand Sm) of the outer surface of a rotating roller for casting employedin the SC method are regulated within a specific range, to therebyimprove uniformity in microstructure of the produced rare earth alloy(Japanese Patent Application Laid-Open (kokai) Nos. 2002-59245 and9-1296). However, the above regulation is carried out in order toprevent change in microstructure in a direction of strip width and toprevent lowering of cooling rate at strip ends. In addition, themorphology of raised/dented segments which provide a surface roughnessis not particularly specified.

In contrast, according to the present invention, change inmicrostructure on the alloy flake in a thickness direction; i.e., fromthe roller side to the free surface side is prevented, to thereby attaina uniform microstructure. The uniformity is determined on the basis ofby fine R-rich phase region, and a specific range of percent volumethereof is provided. In this point, the present invention is completelydifferent from the above inventions (Japanese Patent ApplicationLaid-Open (kokai) Nos. 2002-59245 and 9-1296).

EXAMPLES Example 31

Neodymium, ferroboron, cobalt, aluminum, copper, and iron were mixed tothereby obtain the following alloy composition: Nd: 31.5% by mass; B:1.00% by mass; Co: 1.0% by mass; Al: 0.30% by mass; Cu: 0.10% by mass;and a balance of iron. The resulting mixture was melted in an aluminacrucible in an argon gas atmosphere (1 atm) by use of a high-frequencyinduction melting furnace. The resulting molten alloy was cast throughstrip casting, to thereby prepare alloy flakes.

The rotating roller for casting having a diameter of 300 mm and made ofpure copper was employed. During casting, the inside of the copperroller was cooled by water. The roller had a cast surface roughness, asrepresented by 10-point average roughness (Rz), of 4.0 μm. The surfaceroughness of the cast surface was generally provided by elongatedraised/dented segments which extend in random directions and cross oneanother. The roller was rotated at a peripheral velocity of 1.0 m/s, tothereby produce alloy flakes having a mean thickness of 0.30 mm.

The thus-produced alloy flakes were found to have a surface (mold side)roughness, as represented by 10-point average roughness (Rz), of 4.6 μm.Ten flakes were selected from the alloy flakes and polished in a fixedstate. Each flake was observed under a scanning electron microscope(SEM) and a back-scattered electron image (BEI) was captured at amagnification of ×100. Through analysis of the thus-captured photographby means of an image graphic analyzer, the percent volume of fine R-richphase region was found to be 3% or less.

Comparative Example 31

The procedure of Example 31 including preparing a raw material, melting,and casting through an SC method was repeated, except that a rotatingroller for casting having a surface roughness, as represented by10-point average roughness (Rz), of 4.0 μm was employed. The roller has,on the cast surface, elongated raised/dented segments which extend in arotation direction almost in parallel, and has no substantial segmentswhich cross the above linear segments.

The thus-produced alloy flakes were evaluated in a manner similar tothat of Example 31. The alloy flakes were found to have a surface (moldside) roughness, as represented by 10-point average roughness (Rz), of4.5 μm and to have a percent volume of fine R-rich phase region of 25%.

Comparative Example 32

The procedure of Example 31 including preparing a raw material, melting,and casting through an SC method was repeated, except that a rotatingroller for casting having a surface roughness, as represented by10-point average roughness (Rz), of 100 μm was employed. Similar to thecase of Example 31, the surface roughness was generally provided byelongated raised/dented segments which cross one another.

In Comparative Example 32, a portion of metal remained in contact withthe roller, without coming off the roller, and reached the tundish afterone rotation of the roller. Since the front end of the tundish wasbroken by the alloy, casting operation was stopped.

Working examples of production of sintered magnets will next bedescribed.

Example 32

The alloy flakes produced in Example 31 were subjected to hydrogendecrepitation and micro-pulverization by use of a jet mill. Hydrogenabsorption step—the first step of hydrogen decrepitation—was performedunder the conditions: 100% hydrogen atmosphere, 2 atm, and retentiontime of 1 hour. The temperature of the alloy flakes at the start ofhydrogen absorption reaction was 25° C. Hydrogen desorptionstep—subsequent step—was performed under the conditions: vacuum of 0.133hPa, 500° C., and retention time of 1 hour. To the resultant powder,zinc stearate powder was added in an amount of 0.07% by mass. Themixture was sufficiently mixed in a 100% nitrogen atmosphere by use of aV-blender, and then micro-pulverized by use of a jet mill in a nitrogenatmosphere incorporated with oxygen (4,000 ppm). The resultant powderwas sufficiently mixed again in a 100% nitrogen atmosphere by use of aV-blender. The obtained powder was found to have an oxygen concentrationof 2,500 ppm. Through analysis of the carbon concentration of thepowder, the zinc stearate content of the powder was calculated to be0.05% by mass. The mean particle sizes of the powder, as measured bymeans of a laser diffraction particle size distribution measurementapparatus, were found to be 5.00 μm (D50), 1.98 μm (D10), and 8.51 μm(D90).

Subsequently, the thus-obtained powder was press-molded in a 100%nitrogen atmosphere and a lateral magnetic field by use of a moldingapparatus. The molding pressure was 1.2 t/cm², and the magnetic field inthe mold cavity was controlled to 15 kOe. The thus-obtained compact wasmaintained sequentially in vacuum of 1.33×10^(−5 hPa at) 500° C. for onehour, in vacuum of 1.33×10⁻⁵ hPa at 800° C. for two hours, and in vacuumof 1.33×10⁻⁵ hPa at 1,050° C. for two hours for sintering. The densityof the sintered product was as sufficiently high as 7.5 g/cm³ or more.The sintered product was further heat-treated at 560° C. for one hour inan argon atmosphere, to thereby produce a sintered magnet.

Magnet characteristics of the sintered magnet were measured by means ofa direct-current BH curve tracer. The results are shown in Table 3. Theoxygen content and particle size of the raw micro-powder for producingthe sintered magnet are also shown in Table 3.

Comparative Example 33

In a manner similar to Example 32, alloy flakes produced in ComparativeExample 31 were pulverized, to thereby obtain a micro-powder. Theprocedure of molding and sintering performed in Example 32 was repeated,to thereby produce a sintered magnet.

Magnet characteristics of the sintered magnet produced in ComparativeExample 33 were measured by means of a direct-current BH curve tracer.The results are shown in Table 3. The oxygen content and particle sizeof the raw micro-powder for producing the sintered magnet of ComparativeExample 33 are also shown in Table 3.

TABLE 3 Micro-powder Oxygen Magnet content Particle size (μm) Br iHc(BH)_(max) (ppm) D10 D50 D90 (kG) (kOe) (MGOe) Example 32 2,400 2.0 5.08.5 13.6 14.5 44.5 Comp. Ex. 33 3,000 1.6 5.0 8.8 13.5 13.8 43.5 Example33 — — — —  9.1 13.4 18.0 Comp. Ex. 34 — — — —  9.1 12.7 17.3

As is clear from Table 3, a micro-powder obtained in Comparative Example33 has a smaller D10 as compared with that of the micro-powder obtainedin Example 32; i.e., contains large amounts of very minute particleshaving a particle size of less than about 1 μm. Since such a minutepowder is readily oxidized, the micro-powder obtained in ComparativeExample 33 exhibits a slightly higher oxygen content as compared withthat of the micro-powder of Example 32. Magnetic characteristics of themicro-powder obtained in Comparative Example 33 are inferior to those ofthe micro-powder of Example 32. The poor characteristics are mainlyconsidered to be attributed to an increase in oxygen content and pooruniformity in microcrystal structure.

Working examples of production of bonded magnets will next be described.

Example 33

The procedure of Example 31 including casting through an SC method wasrepeated, except that a raw material having the following alloycomposition: Nd: 28.5%; B: 1.00% by mass; Co: 10.0% by mass; Ga: 0.5% bymass; and a balance of iron was used, to thereby produce alloy flakes.

The thus-produced alloy flakes were evaluated in a manner similar tothat of Example 31. The alloy flakes were found to have a surface (moldside) roughness, as represented by 10-point average roughness (Rz), of4.3 μm and to have a percent volume of fine R-rich phase region of 3% orless. The alloy flakes contain no α-Fe.

The above alloy flakes were subjected to HDDR treatment includingannealing under hydrogen (1 atm) at 820° C. for one hour and subsequentannealing in vacuum at 820° C. for one hour. The resultant alloy powderwas pulverized by means of a Brawn mill so as to have a particle size of150 μm or less and blended with an epoxy resin (2.5% by mass). Theresultant mixture was press-formed in a magnetic field of 1.5 T, tothereby obtain a bonded magnet. Magnetic characteristics of the bondedmagnet are shown in Table 3.

Comparative Example 34

The procedure of Comparative Example 31 including melting and castingthrough an SC method was repeated, except that the raw material wasreplaced by the raw material employed in Example 33, to thereby producealloy flakes. The thus-produced alloy flakes were evaluated in a mannersimilar to that of Example 31. The alloy flakes were found to have asurface (mold side) roughness, as represented by 10-point averageroughness (Rz), of 4.8 μm and to have a percent volume of fine R-richphase region of 30%.

Subsequently, a bonded magnet was produced by use of alloy flakesobtained in Comparative Example 34 in a manner similar to that ofExample 33. Magnetic characteristics of the bonded magnet are shown inTable 3.

As is clear from Table 3, the bonded magnet produced in Example 33exhibits more excellent magnetic characteristics than those of thebonded magnet produced in Comparative Example 34. The bonded magnetproduced in Comparative Example 34 has a high percent volume of fineR-rich phase region and contains a large number of comparatively smallgrains having a grain size of 50 μm or less produced through HDDRtreatment or pulverization. The poor magnetic characteristics areconsidered to be attributable to such a small grain size.

Effects of the Invention

The R-T-B alloy flakes according to the present invention, having asmall percent volume of fine R-rich region, exhibit higher uniformity inR-rich phase dispersion state in the alloy as compared with alloy flakesproduced through a conventional SC method. Thus, sintered magnetsproduced from the R-T-B alloy flakes of the present invention and bondedmagnets produced by use of the flakes through an HDDR method exhibitmore excellent magnetic characteristics than those of conventionalmagnets.

Fourth Aspect

FIG. 7 shows a back-scattered electron image, observed under an SEM(scanning electron microscope), of a cross-section of an Nd-Fe-B alloy(Nd: 31.5 mass %) flake which has been cast through a conventional SCmethod. In FIG. 7, the left side corresponds to the mold side, and theright side to the free surface side.

In FIG. 7, white areas correspond to Nd-rich phase (R-rich phase iscalled Nd-rich phase, since R is Nd). From the center portion to thefree surface side (the surface opposite the mold side) of the alloyflake, the Nd-rich phase assumes the form of lamellar portions extendingin the thickness direction, or the form of a small pool of orientedlamellar fragments. In contrast, the Nd-rich phase on the mold side hasconsiderably minute grains as compared with other portions, and suchgrains are present at random in a region on the mold side. The presentinventor denominates such a region “fine R-rich phase region” (when Rpredominantly comprises Nd, the region is called fine Nd-rich phaseregion) and distinguishes this region from other regions. The fineR-rich phase region is generally formed from the mold side and extendsto the center portion. A portion from the center to the free surfaceside where no fine R-rich phase region is present is called a “normalportion”.

During hydrogen decrepitation of R-T-B alloy flakes for producing asintered magnet, the volume of R-rich phase increases by absorbinghydrogen, thereby forming a fragile hydride. Thus, when hydrogendecrepitation is performed, microcracks are formed along or from theR-rich phase contained in the alloy. In the subsequentmicro-pulverization step, the alloy flakes are crushed by virtue of alarge amount of microcracks generated in hydrogen decrepitation.Therefore, when the R-rich phase is dispersed more finely in the alloy,the particle size of the resultant micro-powder tends to be smaller.Thus, as compared with a normal portion, the fine R-rich phase region isreadily crushed to form minute particles. For example, the alloy powderobtained from a normal portion has an average particle size of about 3μm as measured by means of FSSS (Fisher Sub-Sieve Sizer), whereas thealloy powder obtained from the fine R-rich phase region contains a largeportion of micropowder having a particle size of 1 μm or less, resultingin a broad particle size distribution profile of the micro-pulverizedproduct.

Japanese Unexamined Patent Application, First Publication Nos. 09-170055and 10-36949 disclose that the dispersion state of R-rich phase in anR-T-B alloy can be controlled by regulating a cooling rate of moltenalloy solidified during casting or by heat treatment. However, incontrast to the case of a normal portion, behavior of the R-rich phasepresent in the fine R-rich phase region is difficult to control byregulating a cooling rate of solidified molten metal or by heattreatment, and the R-rich phase is not widely dispersed but remainsfinely dispersed.

The percent volume of the fine R-rich phase region can be determined inthe following manner. FIG. 9 is a back-scattered electron image of thesame observation area as that of FIG. 7, but in FIG. 9 the boundarybetween the fine R-rich phase region and the normal portion is specifiedby the line. Since the boundary between two regions can be readilyidentified through observation of the dispersion state of R-rich phase,the percent area of the fine R-rich phase region in the observation areacan be calculated. The percent area in the cross-section corresponds tothe percent volume of the alloy. The percent area in the cross-sectioncan be determined rapidly and unambiguously by means of a graphic imageanalyzer.

Morphology, size and density of R-rich phase in the fine R-rich phaseregion are different from those of the normal portion. When the fine andgenerally spherical R-rich phase has a certain measure of density, it ispossible to judge it as a fine R-rich phase region. A threshold valuevaries depending on the image quality of the back-scattered electronimage used in graphic image analysis. However, the present inventorshave found the following fact. That is, the portion wherein the densityof R rich phase each having roundness of 1 to 1.4 and area of 5 μm² orless is 20 or more per 100 μm² is judged as a fine R-rich region. Incase the portion, which does not correspond to the standard, exists atthe mold side of the portion which corresponds to the standard, thismold sided portion is also judged as a fine R-rich region. The fineR-rich phase region can be measured with good reproducibility by thismethod. Herein, the term “roundness” refers to a value obtained bydividing (square of the circumference) of the subject figure by(4π×area), and the roundness is 1 in case of circle and increases incase of elongated morphology. Furthermore, the present inventors havefound that the percent volume thus calculated of the fine R-rich phaseregion can successfully describe the feature of the alloy structure andan influence on the particle size distribution profile aftermicro-pulverization and magnetic characteristics of the sintered magnet.

Upon measurement of percent volume of fine R-rich phase region, the fineR-rich phase region content greatly varies among alloy flakes or withinone alloy flake, even when the alloy flakes are cast simultaneously.Thus, graphic image analysis is performed by use of about 5 to about 10flakes and obtained percent area values are averaged, to therebycalculate the percent volume of the fine R-rich phase region for theentirety of the alloy.

FIG. 10 is a back-scattered electron image of a cross-section of anR-T-B alloy flake (Nd: 31.5 mass %) falling within a scope of the FourthAspect of the present invention. In FIG. 10, the left side correspondsto the mold side and the right side to the free surface side. The alloyflake of the Fourth Aspect of the present invention is characterized byemploying a rotating roller for casting, the roller having a surfaceroughness provided by a plurality of elongated raised/dented segmentsformed on the cast surface, almost all of raised/dented segmentsextending in a direction forming at last a specific angle to a rotationdirection of the roller for casting in the method for producing arare-earth-containing alloy flake including a strip casting method. Thealloy flake shown in FIG. 10 is produced by using a rotating roller forcasting, which has the cast surface having the above construction and asurface roughness of 3.5 μm. As shown in FIG. 10, the alloy flake of thepresent invention contains no fine R-rich phase region on the mold side(left side in FIG. 10), and R-rich phase is dispersed, from the moldside to the free surface side, with remarkably excellent uniformity.

In contrast, the alloy flake shown in FIG. 7 is produced by using arotating roller for casting, having, on the cast surface, elongatedraised/dented segments which extend in a rotation direction almost inparallel. The surface roughness of the cast surface is almost the sameas 3.3 μm, however, the percent volume of the fine R-rich phase regionat the left side shown in the drawing is larger than that of the alloyflake shown in FIG. 10.

The relationship between the direction of elongated raised/dentedsegments on the surface of the rotating roller for casting and the fineR-rich phase region of the alloy flake in the strip casting method canbe described as follows.

In case the surface of the rotating roller for casting is smooth and hashigh wettability with respect to the molten alloy, heat is transferredfrom the molten alloy to the mold at remarkably high efficiency (i.e.,heat transfer coefficient is high). Thus, the mold side alloy is rapidlycooled excessively. The fine R-rich phase region is considered to behighly prone to be generated through excessively rapid cooling of theportion of the alloy on the mold side resulting from the large heattransfer coefficient of the molten alloy to the mold.

In contrast, when the surface of the rotating roller for casting isminutely roughened, the surface roughness of the rotating roller forcasting increases. Therefore, the minute irregularities formed on thesurface of the rotating roller for casting cannot be filled completelywith the molten alloy, because of its viscosity. Thus, a portion of thealloy remains not in contact with the roller, thereby lowering the heattransfer coefficient. As a result, a portion of the alloy on the moldside is not rapidly cooled to an excessive extent. Accordingly, theabove mechanism is considered to prevent generation of the fine R-richphase region.

In connection with the morphology of raised/dented segments, when thesesegments are elongated segments, each of contact and non-contactportions between the molten alloy and the roller tends to extend alongan elongated raised/dented segment. Accordingly, the internalmicrostructure is also prone to exhibit continuity along such araised/dented segment. In this case, if a fine R-rich phase region isformed in an elongated raised/dented segment for some reason, therearises the risk of growth of the fine R-rich phase region in the entireportion of the elongated raised/dented segment.

The present inventors have found the following fact. That is, like therotating roller for casting applied to the method of the Fourth aspectof the present invention, when elongated raised/dented segments areformed in a direction intersecting with the rotation direction of therotating roller for casting, the continuity of the internalmicrostructure of the alloy flake along the extending direction ofraised/dented segments can be lowered and also formation of the fineR-rich phase region can be suppressed.

The relationship between the direction of elongated raised/dentedsegments formed on the cast surface and the internal microstructureformed on the alloy flake is considered as follows. That is, in caseelongated raised/dented segments extend in a rotation direction inparallel, or almost in parallel, when the molten alloy is contacted withthe surface of the rotating roller, thereby to fill the space betweenraised/dented segments, an atmospheric gas in the dented portion isextruded to the rotation direction. As a result, the contact areabetween the surface of the roller and the molten alloy increases.

In contrast, in case elongated raised/dented segments extend in adirection intersecting with a roll rotation direction, as the anglebetween the both increases, an atmospheric gas in the dented portion ofthe surface of the rotating roller tends to be trapped with the moltenalloy, thus making it possible to suppress excess contact between thesurface of the roller and the molten alloy. Therefore, even though thesurface roughness of the cast surface is the same, as the angle betweenthe extending direction of elongated raised/dented segments and the rollrotating direction increases, the formation of the fine R-rich phaseregion tends to be suppressed. Also the continuity of the internalmicrostructure along the extending direction of raised/dented segmentsis lowered as compared with the case wherein the rotation directioncorresponds to the extending direction of raised/dented segments.Therefore, if a fine R-rich phase region is formed, there arises therisk of growth of the fine R-rich phase region.

In the Fourth Aspect of the present invention, a rotating roller forcasting is used, the roller being characterized by having, on the castsurface, a plurality of elongated raised/dented segments and having asurface roughness provided by a plurality of elongated raised/dentedsegments, as represented by 10-point average roughness (Rz), fallingwithin a range of 3 μm to 60 μm, 30% or more of raised/dented segmentsamong entire elongated raised/dented segments extending in a directionforming an angle of 30° or more to a roller rotation direction.

Preferably used is a rotating roller for casting, 30% or more ofraised/dented segments among entire elongated raised/dented segmentsextending in a direction forming an angle of 45° or more to a rollerrotation direction, or a rotating roller for casting, 50% or more ofraised/dented segments among entire elongated raised/dented segmentsextending in a direction forming an angle of 30° or more to a rollerrotation direction.

More preferably used is a rotating roller for casting, 50% or more ofraised/dented segments among entire elongated raised/dented segmentsextending in a direction forming an angle of 45° or more to a rollerrotation direction.

Herein, the term “angle between the extending direction of raised/dentedsegments and the roller rotating direction” is defined as 0° in caseelongated raised/dented segments extend in a roller rotation directionin parallel, while it is defined as 90° in case elongated raised/dentedsegments extend in a roller width direction in parallel.

As a result, formation of fine R-rich phase region is prevented, therebyattaining a uniform microstructure, even though the surface roughness issmall compared with the case wherein elongated raised/dented segmentsextend in a rotation direction almost in parallel. Since a small surfaceroughness of the rotating roller for casting decreases the amount ofgrind for regulating the roller surface, the service life of therotating roller for casting can be prolonged. According to the FourthAspect of the present invention, standards for controlling surfaceconditions of the roller can be simplified, since effects exerted bysurface roughness become smaller.

Even when a conventional SC method is employed, the produced alloyflakes include, to some extent, those having a uniform microstructure asshown in FIG. 10. However, alloy flakes having large portions of fineR-rich phase regions as shown in FIG. 7 are also producedsimultaneously, thereby deteriorating uniformity in the entiremicrostructure of the resultant alloy. Failure to attain uniformity inmicrostructure of the alloy produced through a conventional SC methodmay be attributable to difference in conditions of contact between theroller surface and the molten alloy; e.g., the fine surface state of therotating roller for casting, molten alloy supply conditions, and theatmosphere during casting.

Surface irregularity having proper size provided on the surface of arotating roller for casting prevents excessive heat transfer duringsolidification of molten alloy, to thereby suppress, at highreproducibility, generation of fine R-rich phase region. In addition,according to the Fourth Aspect of the present invention, elongatedraised/dented segments on the surface of the roller extend in thedirection forming an angle to a rotation direction of a rotating rollerfor casting. Thus, effect of preventing formation of fine R-rich phaseregion is strengthened and is satisfactory, even when the surfaceroughness is comparatively small. As a result, alloy flakes having sucha uniform microstructure as shown in FIG. 8 can be produced at highyield.

It is not necessarily that elongated raised/dented segments of thepresent invention are continuous, and may be intermittent. Alsoelongated raised/dented segments may in the form of curve or straightline.

The roller is preferably made of pure copper or a copper alloy. Also thecast surface layer can be coated in the present invention.

The present invention will next be described in detail.

(41) Strip Casting Method

The present invention is drawn to a rare-earth-containing alloy flakewhich is produced through the strip casting method. Herein, casting ofR-T-B alloy through the strip casting method will be described.

FIG. 4 is a schematic view showing a casting apparatus employed in stripcasting. Generally, when an R-T-B alloy is cast, the alloy is mademolten by use of a refractory crucible 1 in vacuum or an inert gasatmosphere, because it is highly active. The thus-molten alloy ismaintained at 1,350 to 1,500° C. for a predetermined period of time, andsupplied, via a tundish 2 having optional flow-control means orslag-removing means, to a rotating roller 3 for casting whose interioris cooled with water. The rate of supplying the molten alloy and therotation speed of the rotating roller are appropriately regulated inaccordance with the thickness of the alloy flakes to be produced.Generally, the rotation speed of the rotating roller is about 0.5 toabout 3 m/s (in terms of peripheral velocity). The rotating roller forcasting is preferably made of copper or copper alloy, from the viewpointof high thermal conductivity and availability. The surface of therotating roller for casting is prone to adsorb metallic material,depending on the material and surface conditions of the rotating roller.Thus, provision of an optional cleaning apparatus stabilizes qualitiesof the cast R-T-B alloy. The alloy 4 solidified on the rotating rolleris released from the roller on the side opposite the tundish side andcollected into a collection container 5. The microstructure of R-richphase present in the normal portion can be controlled by means ofheating/cooling means provided in the collection container.

The alloy flake of the present invention preferably has a thickness ofat least 0.1 mm and not greater than 0.5 mm. When the thickness of thealloy flake is less than 0.1 mm, solidification rate increasesexcessively, thereby providing an excessively small crystal grain size,which is equivalent to the particle size of micro-pulverized powderapplied to the magnet production step. In this case, percent orientationand magnetization of the produced magnets are problematicallydeteriorated. A thickness of the alloy flake in excess of 0.5 mm resultsin problems, such as deterioration of R-rich phase dispersibilitystemming from a decrease in solidification rate, and problematicprecipitation of α-Fe.

(42) Surface Roughness of the Cast Surface of the Rotating Roller forCasting

According to the Fourth Aspect of the present invention, when an R-T-Bmagnet alloy is cast through a strip casting method, the surfaceroughness, as represented by 10-point average roughness (Rz), of thecast surface of a rotating roller for casting is controlled to fallwithin a range of 3 μm to 60 μm.

Herein, the term “surface roughness” refers to a surface roughnessdetermined under the conditions specified in JIS B 0601 “Surfaceroughness—Definitions and Designation”, and 10-point average roughness(Rz) is defined therein. Specifically, a surface to be measured is cutwith a plane which is perpendicular thereto, to thereby obtain a contourappearing on a cut end (profile curve). Any surface waviness componentlonger than a prescribed wavelength is cut off from the profile curve bymeans of a phase-compensation-type high-pass filter or a similar device,to thereby obtain a curve (roughness curve). Only the reference lengthis sampled from the roughness curve in the direction of its mean line,and the sum of the average value of absolute values of the heights ofthe five highest profile peaks (Yp) and the depths of the five deepestprofile valleys (Yv) measured in the vertical direction from the meanline of this sampled portion is calculated, to thereby obtain the10-point average roughness (Rz). Measurement parameters such asreference length are defined in the above JIS B 0601, as standard valuesof reference length for determining corresponding surface roughnessvalues.

Since the surface roughness often varies in a wide range among samplesto be measured, an average value of surface roughness for at least fiveflakes should be employed.

(43) Morphology of Surface Irregularity of the Cast Surface of a Roller

According to the Fourth Aspect of the present invention, surfaceirregularities are generally provided by a plurality of elongatedraised/dented segments formed on the cast surface, almost all ofelongated raised/dented segments being formed while extending in adirection forming at least a specific angle to a rotation direction ofthe roller for casting. Specifically, 30% or more of raised/dentedsegments extend in a direction forming an angle of 30° or more to aroller rotation direction. Preferably, 30% or more of raised/dentedsegments extend in a direction forming an angle of 45° or more to aroller rotation direction, or 50% or more of raised/dented segmentsamong entire elongated raised/dented segments extend in a directionforming an angle of 30° or more to a roller rotation direction. Morepreferably, 50% or more of raised/dented segments among entire elongatedraised/dented segments extend in a direction forming an angle of 45° ormore to a roller rotation direction.

In case elongated raised/dented segments extend in a rotation directionin parallel, or almost in parallel, when the molten alloy is contactedwith the surface of the rotating roller, thereby to fill the spacebetween raised/dented segments, an atmospheric gas in the dented portionis likely to be extruded. As a result, the contact area between thesurface of the roller and the molten alloy increases. However, as theangle between the extending direction of raised/dented segments and therotation direction increases, an atmospheric gas in the dented portionof the surface of the rotating roller tends to be trapped with themolten alloy, thus making it possible to suppress excess contact betweenthe surface of the roller and the molten alloy. Therefore, even thoughthe surface roughness of the cast surface is the same, as the anglebetween the extending direction of elongated raised/dented segments andthe roll rotating direction increases or as the number of elongatedraised/dented segments extending in a direction forming a large angle tothe roller rotation direction, the formation of the fine R-rich phaseregion tends to be suppressed. Also the continuity of the internalmicrostructure along the extending direction of raised/dented segmentsis lowered as compared with the case wherein the rotation directioncorresponds to the extending direction of raised/dented segments.Therefore, if a fine R-rich phase region is formed, there arises therisk of growth of the fine R-rich phase region.

A plurality of elongated raised/dented segments formed on the castsurface can exert the effect of suppressing the above-described fineR-rich phase region when 30% or more of them extend in a directionforming an angle of 30° or more to a roller rotation direction. Noeffect is exerted when the proportion is 30% or less.

It is not necessarily that elongated raised/dented segments on thesurface of the rotating roller for casting of the present invention arecontinuous, and may be intermittent. Also elongated raised/dentedsegments may in the form of curve or straight line.

These elongated raised/dented segments can be formed by forming an anglebetween the polishing direction and a rotation direction even when usingan apparatus equipped with a rotating abrasive paper, a rotary wirebrush, or a belt abrasive apparatus equipped with an abrasive paperwhich linearly transfers.

According to the Fourth Aspect of the present invention, uniformmicrostructure can be provided through effect of elongated raised/dentedsegments extending in an direction forming an angle to a rotationdirection of a roller for casting, even though the surface roughness iscomparatively small.

However, when the surface roughness is 3 μm or less, effect exerted bythe presence of irregularities of the surface of a rotating roller forcasting is unsatisfactory. Thus, heat transfer is promoted throughincreased contact between the molten alloy and the surface of a rotatingroller for casting, thereby readily forming fine R-rich phase region inthe alloy.

When the surface roughness of the rotating roller for casting is inexcess of 60 μm, a solidified alloy flake is engaged with the rollersurface and difficult to peel from the roller, thereby possibly causingtrouble such as breakage of a tundish. Therefore, the surface roughnessof the rotating roller for casting is controlled to 60 μm or less.

(44) Percent Volume of Fine R-Rich Phase Region in the Alloy

In case an R-T-B alloy is produced by the method of the presentinvention, the percent volume of fine R-rich phase region in the R-T-Balloy is regulated to 20% or less. Thus, the alloy powder which has beenmicro-pulverized for producing sintered magnets has a sharp particlesize distribution profile, thereby yielding sintered magnets withoutvariation in characteristics.

(45) Method for Producing Rare Earth Sintered Magnet Alloy Powder andMethod for Producing Rare Earth Sintered Magnets

The rare earth magnet alloy flakes formed of R-T-B alloy for producing amagnet which flakes have been cast through the method according to thepresent invention are pulverized, shaped, and sintered, to therebyproduce anisotropic sintered magnets of excellent characteristics.

Typically, pulverization of the alloy flakes is sequentially performedin the order of hydrogen decrepitation and micro-pulverization, tothereby produce an alloy powder having a size of approximately 3 μm(FSSS).

In the present invention, hydrogen decrepitation includes a hydrogenabsorption step as a first step and a hydrogen desorption step as asecond step. In the hydrogen absorption step, hydrogen is caused to beabsorbed predominantly in the R-rich phase of alloy flakes in a hydrogengas atmosphere at 266 hPa to 0.3 MPa. The R-rich phase is expanded involume due to R hydride generated in this step, to thereby minutelybreak the alloy flakes themselves or generate numerous micro-cracks.Hydrogen absorption is carried out within a temperature range of ambienttemperature to approximately 600° C. However, in order to increaseexpansion in volume of R-rich phase so as to effectively reduce theflakes in size, hydrogen absorption is preferably performed underincreased hydrogen gas pressure and within a temperature range ofambient temperature to approximately 100° C. The time for hydrogenabsorption is preferably one hour or longer. The R hydride formedthrough the hydrogen absorption step is unstable in the atmosphere andreadily oxidized. Thus, the hydrogen-absorbed product is preferablysubjected to hydrogen desorption treatment by maintaining the alloyflakes at about 200 to about 600° C. in vacuum of 1.33 hPa or less.Through this treatment, R hydride can be transformed into a productstable in the atmosphere. The time for hydrogen desorption treatment ispreferably 30 minutes or longer. If the atmosphere is controlled forpreventing oxidation during steps to be carried out after hydrogenabsorption to sintering, hydrogen desorption treatment can also beomitted.

The R-T-B alloy flake produced through the strip casting methodaccording to the present invention is characterized in that R-rich phaseis uniformly dispersed in the alloy flake. The average inter R-richphase spacing, which depends on the particle size of the pulverizedpowder for producing magnets, is preferably 3 μm to 8 μm. Duringhydrogen decrepitation, cracks are introduced to the alloy flake alongor from the R-rich phase therein. Therefore, micro-pulverization of aproduct which has undergone hydrogen decrepitation attains, to a maximumdegree, the effect of the R-rich phase uniformly and finely dispersed inthe alloy, thereby effectively producing an alloy powder exhibiting aremarkably sharp particle size distribution profile. When sinteredmagnets are produced without performing the hydrogen decrepitation step,the produced sintered magnets have poor characteristics (M. Sagawa etal., Proceeding of the 5th international conference on Advancedmaterials, Beijing, China (1999)).

Micro-pulverization is a step of pulverizing R-T-B alloy flakes forattaining a particle size of approximately 3 μm (FSSS). Amongpulverizers for performing the micro-pulverization, a jet mill is mostpreferred, in view of high productivity and a sharp particle sizedistribution profile. By use of alloy flakes according to the presentinvention having a low fine R-rich phase region content, an alloy powderexhibiting a sharp particle size distribution profile can be produced athigh efficiency without variation.

Upon micro-pulverization, the atmosphere is controlled to an inert gasatmosphere such as an argon gas atmosphere or nitrogen gas atmosphere.The inert gas may contain oxygen in an amount of 2% by mass or less,preferably 1% by mass or less. The presence of oxygen enhancespulverization efficiency and attains oxygen concentration of the alloypowder produced through pulverization to 1,000 to 10,000 ppm, to therebyappropriately stabilize the alloy powder. In addition, abnormal graingrowth during sintering to form magnets can be prevented.

When the alloy powder is molded in a magnetic field, in order to reducefriction between the powder and the inner wall of a mold and to reducefriction generated among powder particles for enhancing orientation, alubricant such as zinc stearate is preferably added to the powder. Theamount of the lubricant to be added is 0.01 to 1% by mass. Although thelubricant may be added before or after micro-pulverization, thelubricant is preferably mixed sufficiently, before molding in magneticfield, in an inert gas atmosphere such as argon gas or nitrogen gas byuse of a mixing apparatus such as a V-blender.

The R-T-B alloy powder having a particle size of about 3 μm (FSSS)obtained through micro-pulverization is press-molded in magnetic fieldby use of a molding apparatus. The mold to be employed is fabricatedfrom a magnetic material and a non-magnetic material in combination inconsideration of the orientation of magnetic field in the mold cavity.The pressure at molding is preferably 0.5 to 2 t/cm², and the magneticfield in the mold cavity during molding is preferably 5 to 20 kOe. Theatmosphere during molding is preferably an inert gas atmosphere such asargon gas or nitrogen gas. However, if the powder has been subjected tothe aforementioned anti-oxidation treatment, molding can be performed inair.

Molding may be performed through cold isostatic pressing (CIP) or rubberisostatic pressing (RIP) employing a rubber mold. Since the alloy powderis pressed isostatically through CIP or RIP, variation in orientation ofmagnetization during press-molding is lowered. Thus, the degree oforientation of the produced compact can be increased as compared withthat produced by use of a metal mold, and maximum magnetic energyproduct can be enhanced.

Sintering of the compact is performed at 1,000 to 1,100° C. Theatmosphere during sintering is preferably an argon gas atmosphere or avacuum atmosphere of 1.33×10⁻² hPa or less. A retention time at thesintering temperature of one hour or longer is preferred. Duringsintering, prior to reaching the sintering temperature, a lubricant andhydrogen must be removed as completely as possible from a compact to besintered. The lubricant is removed by maintaining the compact preferablyunder the conditions: in vacuum of 1.33×10⁻² hPa or less or under an Arflow atmosphere at reduced pressure; at 300 to 500° C.; and for 30minutes or longer. Hydrogen is removed by maintaining the compactpreferably under the conditions: in vacuum of 1.33×10⁻² hPa or less; at700 to 900° C.; and for 30 minutes or longer.

After completion of sintering, in order to enhance the coercivity ofsintered magnet to be produced, the sintered product may be treated at500 to 650° C. in accordance with needs. An argon gas atmosphere or avacuum atmosphere is preferred, and a retention time of 30 minutes orlonger is preferred.

The rare earth magnet R-T-B alloy flake produced through the methodaccording to the present invention in which formation of fine R-richregion is suppressed can be used suitably for producing bonded magnetsas well as sintered magnets. Production of a bonded magnet by use of therare earth magnet alloy flakes according to the present invention willnext be described.

Firstly, the R-T-B alloy flakes of the present invention undergo heattreatment in advance in accordance with needs. The heat treatment isperformed in order to remove α-Fe contained in the alloy and to coarsencrystal grains. The production steps of the alloy powder for producingbonded magnets includehydrogenation-disproportionation-desorption-recombination (HDDR)treatment. However, α-Fe present in the alloy cannot be removed in theHDDR treatment step, and remaining α-Fe deteriorates magnetism.Therefore, α-Fe must be removed prior to performing the HDDR treatment.

The alloy powder for producing bonded magnets has a mean particle sizeof 50 to 300 μm, which is considerably greater than that of the alloypowder for producing sintered magnets. When the bonded magnet alloyflakes undergo HDDR treatment, crystal orientation of recombined crystalgrains of sub-micron size coincides with crystal orientation of crystalgrains of the starting alloy flakes with a certain range of variance.Thus, when two or more crystal grains having different crystalorientations are contained in each of starting alloy flakes, eachparticle of the bonded magnet alloy powder produced from such alloyflakes will contain crystal grains having different crystalorientations. Thus, the alloy powder includes regions having greatvariance in crystal orientation. In such region, the degree oforientation deteriorates, and maximum magnetic energy product of themagnet is low. In order to avoid such deterioration, the crystal grainscontained in the alloy flakes preferably have a large grain size. Thealloy cast through a rapid-cooling/solidification method (e.g., stripcasting) is prone to have a comparatively small crystal grain size.Thus, coarsening of crystal grains through heat treatment is effectivefor enhancing magnet characteristics.

There are many reports in connection with the method for producing abonded magnet alloy powder through the HDDR method (e.g., T. Takeshitaet al., Proc. 10th Int. Workshop on RE magnets and their application,Kyoto, Vol. 1, P. 551 (1989)). Production of the alloy powder throughthe HDDR method is performed in the following manner.

When R-T-B alloy flakes serving as raw material are heated in a hydrogenatmosphere, the R₂T₁₄B phase, a magnetic phase, decomposes at about 700°C. to about 850° C., to thereby form three phases; i.e., α-Fe, RH₂, andFe₂B. Subsequently, in order to remove hydrogen, the hydrogen atmosphereis replaced by an inert gas atmosphere or a vacuum atmosphere, and thetemperature is maintained approximately in the above range. As a result,separated phases are recombined, to thereby form the R₂T₁₄B phase havingan approximately sub-micron crystal grain size. Upon the above process,if the composition of the alloy or treatment conditions areappropriately modified, the magnetization-easy axis of each recombinedR₂T₁₄B phase (C-axis of R₂T₁₄B phase) is aligned approximately inparallel to the C-axis of R₂T₁₄B phase present in the raw material alloybefore decomposition. Thus, there can be produced an anisotropic magnetpowder in which the magnetization-easy axis of minute crystal grains isaligned.

The alloy which has undergone HDDR treatment is pulverized to form analloy powder having a particle size of about 50 to about 300 μm. By useof the alloy powder, a bonded magnet is produced through a processincluding mixing with resin and press-molding or injection-molding.

Similar to the case of the aforementioned hydrogen decrepitation, fineR-rich phase region is prone to form a micro-powder through HDDRtreatment. Characteristics of the magnetic powder obtained through aHDDR method are deteriorated, as the particle size thereof decreases.Thus, the R-T-B alloy of the present invention in which formation offine R-rich phase is suppressed is suitably used in production a bondedmagnet powder including HDDR treatment.

Recently, it has been reported that surface roughness parameters (Sm/Raand Sm) of the outer surface of a rotating roller for casting employedare regulated within a specific range, to thereby improve uniformity inmicrostructure of the produced rare earth alloy (Japanese PatentApplication Laid-Open (kokai) Nos. 2002-59245 and 9-1296). However, theabove regulation is carried out in order to prevent change inmicrostructure in a direction of strip width and to prevent lowering ofcooling rate at strip ends. In addition, the morphology of raised/dentedsegments which provide a surface roughness is not particularlyspecified.

In contrast, according to the present invention, change inmicrostructure on the alloy flake in a thickness direction; i.e., fromthe roller side to the free surface side is prevented, to thereby attaina uniform microstructure. The uniformity is determined on the basis ofby fine R-rich phase region, and a specific range of percent volumethereof is provided. In this point, the present invention is completelydifferent from the above inventions (Japanese Patent ApplicationLaid-Open (kokai) Nos. 2002-59245 and 9-1296).

Japanese Patent Application Laid-Open (kokai) No. 4-28457 proposes toreduce a change in grain size by regulating the roughness Ra of thesurface of the roller within a range of 0.05 μm to 1.5 μm. Although theinvention in the publication relates to a rapid quenching having a muchhigher roller speed and the standard for the measurement of the surfaceroughness is different from that in the present invention, the presentinventors have found that Ra of the surface of the roller for casting ofthe present invention is larger than that of the above invention.

EXAMPLES Example 41

Neodymium, ferroboron, cobalt, aluminum, copper, and iron were mixed tothereby obtain the following alloy composition: Nd: 31.5% by mass; B:1.00% by mass; Co: 1.0% by mass; Al: 0.30% by mass; Cu: 0.10% by mass;and a balance of iron. The resulting mixture was melted in an aluminacrucible in an argon gas atmosphere (1 atm) by use of a high-frequencyinduction melting furnace. The resulting molten alloy was cast throughstrip casting, to thereby prepare alloy flakes. The rotating roller forcasting having a diameter of 300 mm and made of pure copper wasemployed. During casting, the inside of the copper roller was cooled bywater. The roller had a cast surface roughness, as represented by10-point average roughness (Rz), of 4.0 μm. The surface roughness of thecast surface was generally provided by elongated raised/dented segmentsformed on the cast surface, 50% or more of raised/dented segmentsextending in a direction forming an angle of 45° or more to a rollerrotation direction. The roller was rotated at a peripheral velocity of1.0 m/s, to thereby produce alloy flakes having a mean thickness of 0.30mm.

Ten flakes were selected from the resulting alloy flakes and polished ina fixed state. Each flake was observed under a scanning electronmicroscope (SEM) and a back-scattered electron image (BEI) was capturedat a magnification of ×100. Through analysis of the thus-capturedphotograph by means of an image graphic analyzer, the percent volume offine R-rich phase region was found to be 3% or less.

Comparative Example 41

The procedure of Example 41 including preparing a raw material, melting,and casting through an SC method was repeated, except that a rotatingroller for casting having a surface roughness, as represented by10-point average roughness (Rz), of 4.0 μm was employed. The roller has,on the cast surface, elongated raised/dented segments which extend in arotation direction almost in parallel, and has no substantial segmentswhich incline to the rotation direction.

The thus-produced alloy flakes were evaluated in a manner similar tothat of Example 1. The alloy flakes have a percent volume of fine R-richphase region of 25%.

Comparative Example 42

The procedure of Example 41 including preparing a raw material, melting,and casting through an SC method was repeated, except that a rotatingroller for casting having a surface roughness, as represented by10-point average roughness (Rz), of 100 μm was employed. Similar to thecase of Example 1, the surface roughness was generally provided byelongated raised/dented segments formed on the cast surface, 50% or moreof raised/dented segments extending in a direction forming an angle of45° or more to a roller rotation direction.

In Comparative Example 42, a portion of metal remained in contact withthe roller, without coming off the roller, and reached the tundish afterone rotation of the roller. Since the front end of the tundish wasbroken by the alloy, casting operation was stopped.

Working examples of production of sintered magnets will next bedescribed.

Example 42

The alloy flakes produced in Example 41 were subjected to hydrogendecrepitation and micro-pulverization by use of a jet mill. Hydrogenabsorption step, the step preceding hydrogen decrepitation, wasperformed under the conditions: 100% hydrogen atmosphere, 2 atm, andretention time of 1 hour. The temperature of the alloy flakes at thestart of hydrogen absorption reaction was 25° C. Hydrogen desorptionstep, subsequent step, was performed under the conditions: vacuum of0.133 hPa, 500° C., and retention time of 1 hour. To the resultantpowder, zinc stearate powder was added in an amount of 0.07% by mass.The mixture was sufficiently mixed in a 100% nitrogen atmosphere by useof a V-blender, and then micro-pulverized by use of a jet mill in anitrogen atmosphere incorporated with oxygen (4,000 ppm). The resultantpowder was sufficiently mixed again in a 100% nitrogen atmosphere by useof a V-blender. The obtained powder was found to have an oxygenconcentration of 2,500 ppm. Through analysis of the carbon concentrationof the powder, the zinc stearate content of the powder was calculated tobe 0.05% by mass. The mean particle sizes of the powder, as measured bymeans of a laser diffraction particle size distribution measurementapparatus, were found to be 5.11 μm (D50), 1.90 μm (D10), and 8.60 μm(D90).

Subsequently, the thus-obtained powder was press-molded in a 100%nitrogen atmosphere and a lateral magnetic field by use of a moldingapparatus. The molding pressure was 1.2 t/cm², and the magnetic field inthe mold cavity was controlled to 15 kOe. The thus-obtained compact wasmaintained sequentially in vacuum of 1.33×10⁻⁵ hPa at 500° C. for onehour, in vacuum of 1.33×10⁻⁵ hPa at 800° C. for two hours, and in vacuumof 1.33×10⁻⁵ hPa at 1,050° C. for two hours for sintering. The densityof the sintered product was as sufficiently high as 7.5 g/cm³ or more.The sintered product was further heat-treated at 560° C. for one hour inan argon atmosphere, to thereby produce a sintered magnet.

Magnet characteristics of the sintered magnet were measured by means ofa direct-current BH curve tracer. The results are shown in Table 4. Theoxygen content and particle size of the raw micro-powder for producingthe sintered magnet are also shown in Table 4.

Comparative Example 43

In a manner similar to Example 42, alloy flakes produced in ComparativeExample 41 were pulverized, to thereby obtain a micro-powder. Theprocedure of molding and sintering performed in Example 2 was repeated,to thereby produce a sintered magnet.

Magnet characteristics of the sintered magnet produced in ComparativeExample 43 were measured by means of a direct-current BH curve tracer.The results are shown in Table 4. The oxygen content and particle sizeof the raw micro-powder for producing the sintered magnet of ComparativeExample 43 are also shown in Table 4.

TABLE 4 Micro-powder Oxygen Magnet content Particle size (μm) Br iHc(BH)_(max) (ppm) D10 D50 D90 (kG) (kOe) (MGOe) Example 42 2,500 1.9 5.18.6 13.6 14.5 44.4 Comp. Ex. 43 3,000 1.6 5.0 8.8 13.5 13.8 43.5 Example43 — — — —  9.1 13.5 18.0 Comp. Ex. 44 — — — —  9.1 12.7 17.3

As is clear from Table 4, a micro-powder obtained in Comparative Example43 has a smaller D10 as compared with that of the micro-powder obtainedin Example 42; i.e., contains large amounts of very minute particleshaving a particle size of less than about 1 μm. Since such a minutepowder is readily oxidized, the micro-powder obtained in ComparativeExample 43 exhibits a slightly higher oxygen content as compared withthat of the micro-powder of Example 42. Magnetic characteristics of themicro-powder obtained in Comparative Example 43 are inferior to those ofthe micro-powder of Example 42. The poor characteristics are mainlyconsidered to be attributed to an increase in oxygen content and pooruniformity in microcrystal structure.

Working examples of production of bonded magnets will next be described.

Example 43

The procedure of Example 41 including casting through an SC method wasrepeated, except that a raw material having the following alloycomposition: Nd: 28.5%; B: 1.00% by mass; Co: 10.0% by mass; Ga: 0.5% bymass; and a balance of iron was used, to thereby produce alloy flakes.

The thus-produced alloy flakes were evaluated in a manner similar tothat of Example 41. The alloy flakes were found to have a percent volumeof fine R-rich phase region of 3% or less. The alloy flakes contain noα-Fe.

The above alloy flakes were subjected to HDDR treatment includingannealing under hydrogen (1 atm) at 820° C. for one hour and subsequentannealing in vacuum at 820° C. for one hour. The resultant alloy powderwas pulverized by means of a Brawn mill so as to have a particle size of150 μm or less and blended with an epoxy resin (2.5% by mass). Theresultant mixture was press-formed in a magnetic field of 1.5 T, tothereby obtain a bonded magnet. Magnetic characteristics of the bondedmagnet are shown in Table 4.

Comparative Example 44

The procedure of Comparative Example 41 including melting and castingthrough an SC method was repeated, except that the raw material wasreplaced by the raw material employed in Example 3, to thereby producealloy flakes. The thus-produced alloy flakes were evaluated in a mannersimilar to that of Example 41. The alloy flakes were found to have apercent volume of fine R-rich phase region of 30%.

Subsequently, a bonded magnet was produced by use of alloy flakesobtained in Comparative Example 44 in a manner similar to that ofExample 43. Magnetic characteristics of the bonded magnet are shown inTable 4.

As is clear from Table 4, the bonded magnet produced in Example 43exhibits more excellent magnetic characteristics than those of thebonded magnet produced in Comparative Example 44. The bonded magnetproduced in Comparative Example 44 has a high percent volume of fineR-rich phase region and contains a large number of comparatively smallgrains having a grain size of 50 μm or less produced through HDDRtreatment or pulverization. The poor magnetic characteristics areconsidered to be attributable to such a small grain size.

Consequently, according to the method for producing arare-earth-containing alloy flake of the present invention, there can beproduced a rare-earth-containing alloy flake having a small percentvolume of fine R-rich region, exhibit higher uniformity in R-rich phasedispersion state in the alloy as compared with alloy flakes producedthrough a conventional SC method. Thus, sintered magnets produced fromthe resulting rare-earth-containing alloy flake and bonded magnetsproduced by use of the flakes through an HDDR method exhibit moreexcellent magnetic characteristics than those of conventional magnets.

INDUSTRIAL APPLICABILITY

The present invention relates to an improvement of the magneticproperties of rare earth magnets, which are used for various magneticmediums such as hard discs, MRI (Magnetic resonance Imaging), andmotors. The improvement relates to the composition of the rare earthmagnet, method of producing the flakes used for the material of themagnets, and methods of fabricating the solid magnets.

1. A main phase alloy for a rare earth magnet to be processed through atwo-alloy-blending method, the main phase alloy containing R (Rrepresents at least one rare earth element including Y) in an amount of26 to 30% by mass and B in an amount of 0.9 to 1.1% by mass, with thebalance being T (T represents transition metals including Fe as anessential element), characterized in that R has a Pr content of at least5% by mass and the main phase alloy has a percent volume of regioncontaining α-Fe on the basis of the entire microstructure of 5% or less,and wherein at least one surface thereof has a surface roughness, asrepresented by 10-point average roughness (Rz), falling within a rangeof 5 μm to 50 μm.
 2. A main phase alloy for a rare earth magnetaccording to claim 1, wherein R has a Pr content of at least 15% bymass.
 3. A main phase alloy for a rare earth magnet according to claim2, wherein R has a Pr content of at least 30% by mass.
 4. A main phasealloy for a rare earth magnet according to claim 1, wherein at least onesurface thereof has a surface roughness, as represented by 10-pointaverage roughness (Rz), falling within a range of 7 μm to 25 μm.
 5. Amethod for producing a main phase alloy for a rare earth magnetaccording to claim 1, wherein the method comprises strip casting.
 6. Amethod for producing a main phase alloy for a rare earth magnetaccording to claim 5, wherein the surface roughness, as represented by10-point average roughness (Rz), of the cast surface of a rotatingroller for casting is adjusted to fall within a range of 5 μm to 100 μm.7. A method for producing a main phase alloy for a rare earth magnetaccording to claim 5, wherein the surface roughness, as represented by10-point average roughness (Rz), of the cast surface of a rotatingroller for casting is adjusted to fall within a range of 10 μm to 50 μm.8. A method for producing a main phase alloy for a rare earth magnet tobe processed through a two-alloy-blending method, the main phase alloycontaining R (R represents at least one rare earth element including Y)in an amount of 26 to 30% by mass and B in an amount of 0.9 to 1.1% bymass, with the balance being T (T represents transition metals includingFe as an essential element), wherein R has a Pr content of at least 5%by mass and the main phase alloy has a percent volume of regioncontaining α-Fe on the basis of the entire microstructure of 5% or less,and wherein at least one surface thereof has a surface roughness, asrepresented by 10-point average roughness (Rz), falling within a rangeof 5 μm to 50 μm, the process comprising a centrifugal casting methodincluding depositing and solidifying a molten metal on an inner surfaceof a cylindrical mold which is rotating.
 9. A mixed powder for a rareearth sintered magnet produced by mixing a main phase alloy for a rareearth magnet according to any one of claims 1 to 3 with a boundary phasealloy which has an R content higher than that of the main phase alloyand has a Pr content of R lower than that of the main phase alloy.
 10. Amixed powder for a rare earth sintered magnet according to claim 9,wherein the boundary phase alloy contains substantially no Pr.
 11. Arare earth sintered magnet produced through a powder metallurgicalmethod making use of a mixed powder for a rare earth magnet according toclaim 9.